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Combined Fe and O effects on microstructural evolution and strengthening in Cu–Fe nanocrystalline alloys
Materials Science & Engineering A 772 (2020) 138800
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Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea
Combined Fe and O effects on microstructural evolution and strengthening in CuFe nanocrystalline alloys Jinming Guo a, b, *, Qinqin Shao a, c, Oliver Renk d, Lei Li b, Yunbin He b, Zaoli Zhang a, d, Reinhard Pippan a a
Erich Schmid Institute of Materials Science, Austrian Academy of Sciences, 8700 Leoben, Austria School of Materials Science and Engineering, Hubei University, 430062 Wuhan, China Center for High Resolution Electron Microscopy, College of Materials Science and Engineering, Hunan University, 410082 Changsha, China d Department of Materials Science, Chair of Materials Physics, Montanuniversit€ at Leoben, 8700 Leoben, Austria b c
A R T I C L E I N F O
A B S T R A C T
Keywords: Nanocrystalline alloy Severe plastic deformation High pressure torsion Transmission electron microscopy Oxygen effect Microstructural evolution
Immiscible Cu–Fe composites with various compositions are mechanically alloyed by means of high pressure torsion directly from blended powders and vacuum arc-melted bulk materials respectively, which contain different contents of oxygen. All investigated compositions are deformed to extremely large strains reaching a saturated state consisting only of a single face-centered cubic structure. It is found that Fe alloying as well as O addition reduce the achievable grain sizes. The single phase supersaturated solid solutions show a reduced hardness compared to the Cu and Fe heterostructures with similar grain sizes, reflected in a decreasing hardness with increasing strain. In addition, oxygen introduced during powder processing has a significant influence on the microstructures, defect densities and thus properties. Samples generated from powders have smaller grain sizes and much higher dislocation and twin densities than the corresponding arc-melted bulk samples with the same nominal composition. The combined effects of Fe and O on the deformation behavior and hardness of Cu–Fe nanocrystalline alloys are discussed. The current work shows that microstructure and properties can be tuned not only by alloying, but even further by incorporating oxygen, which may offer another design strategy for nanocrystalline alloys.
1. Introduction Severe Plastic Deformation (SPD) offers many advantages in pro ducing bulk nanocrystalline alloys, i. e., dissolution of second phases in usually “immiscible” systems [1], considerable grain refinement down to the nanometer range [2,3]. In addition, via specifically designing and controlling nano-scaled features induced by SPD, such as deformation twins, non-equilibrium grain boundaries, dislocation substructures, so lute segregation and clustering [4–6], many nanocrystalline materials show extraordinary properties compared to their coarse-grained coun terparts [7]. Over the past decades, Cu-based nanocrystalline alloys have drawn intensive interests because of their superior strength compared to coarse-grained pure Cu. They hold a great potential for various applications, such as welding electrodes, heat sinks and elec trical applications due to their excellent electrical and thermal con ductivity in combination with high strength, wear and corrosion resistances [8]. In our previous work, we have successfully fabricated
Cu–Cr [1,3] and Cu–Fe [9,10] nanocrystalline alloys through high pressure torsion (HPT) technique, which is one of the most convenient method of SPD and allows to apply extremely large strains to bulk ma terials. For example, we refined the coarse-grained Cu–Cr alloys down to a saturated grain size of about 15 nm by deformation up to a strain of 340. Finally, in contrast to the conventionally immiscible Cu–Cr system under equilibrium conditions, a maximum amount of 32 at.% Cu was fully dissolved into the Cr matrix forming a stable body-centered cubic structure [1,3]. However, one prominent problem resulted from the consolidation and severe deformation of powder mixtures by SPD is the unavoidable introduction of oxygen [11]. The intrusive oxygen usually can enhance hardness by stabilizing finer grain sizes, which commonly has a negative effect on ductility, causing a premature fracture of the sample even in the elastic regime [12,13]. In addition, our previous work revealed the atomic-scale formation processes of oxides during heating of oxygen-containing nanocrystalline alloys [9].
* Corresponding author. Erich Schmid Institute of Materials Science, Austrian Academy of Sciences, 8700 Leoben, Austria. E-mail addresses: [email protected], [email protected] (J. Guo). https://doi.org/10.1016/j.msea.2019.138800 Received 26 September 2019; Received in revised form 7 December 2019; Accepted 9 December 2019 Available online 10 December 2019 0921-5093/© 2019 Elsevier B.V. All rights reserved.
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Therefore, understanding the oxygen effect in severely deformed bulk nanocrystalline alloys and its synergetic mechanism with other alloying elements is crucial for the future design of high performance materials. In this work, a series of Cu–Fe alloys having different Fe contents are synthesized by HPT until a saturated state consisting only of a single face-centered cubic (fcc) phase is achieved. Two different types of initial materials, blended powders and arc-melted bulk, were inves tigated to understand the oxygen effect. Effects of Fe alloying and O addition on the resulting microstructures, valence states, defects den sities and mechanical properties are revealed via systematic comparison between the different types of samples based on massive statistics. This work poses a promising route to manipulate mechanical properties of nanocrystalline alloys by comprehensive adjustments of solute alloying and intentional oxygen incorporation.
investigated by X-ray Photoelectron Spectroscopy (XPS, ESCALAB 250Xi, Thermo Fisher Scientific, Waltham, USA), employing Al Kα ra diation with a beam size of 500 μm. All samples are fully polished with a media of ethyl alcohol and then transferred to XPS chamber immedi ately, which is followed by Ar ion sputtering with ion energy of 3 KeV for 5 min to completely remove the possible surface oxide layers. Vickers microhardness measurements are conducted on a Buehler Mircomet 5100 using a load of 300 g (HV0.3). Indents are made across the radii of the disks with a spacing of 300 μm between the indents, and average values of six individual measurements on deformed disks for each deformation condition are reported in this paper. Edge dislocation and twin densities are calculated based on statistics of at least 80 high-resolution transmission electron microscopy (HRTEM) images. Edge dislocation density is obtained via dividing the number of dissociated dislocation component 2a < 100 > in HRTEM images taken along [001] zone axes by the total area of the corre sponding grains. Similarly, twin density is calculated via dividing the total length of twin boundaries in HRTEM images taken along [011] zone axes by the total area of the corresponding grains.
2. Experimental procedures The raw materials with nominal compositions of (100–x) at.%Cu – x at.%Fe (x ¼ 0, 5, 15, 25, 35, abbreviated as (100–x)Cu–xFe) are mixed from Cu and Fe powders with purity of 99.9þ% and generated by arcmelting process using Cu and Fe ingots with purity of 99.95%. Here it should be emphasized that a two-step HPT processing technique is used to generate the powder samples, and details of this technique can be found in Ref. [14]. The initial materials for powder samples mentioned in this work refer to the 8 mm disks extracted from the large 40 mm consolidated and HPT deformed disks (1st HPT step with strain of 10 – 54) as shown in the inset of Fig. 1a. The as-generated initial materials are then HPT deformed at room temperature accompanied by air cooling during the whole process. The pressure during HPT is fixed at 7.4 GPa and the rotation speed is 0.4 rotation/min for all samples. All samples are deformed into disks with radii of 4 mm. As in HPT the applied strain is a function of the disk radius, the data shown in this work is either presented as a function of strain ε or given for a certain strain (at radius of 3.0 mm from the disk center). The applied strain ε is calculated based on equation (1) shown in our previous work [3]. For powder samples, the effective strain represents only the strain applied during second-step deformation. XRD is conducted on all samples using Smartlab X-Ray Diffractom eter (Rigaku, Japan) with Cu Kα1 radiation (λ ¼ 1.540593 Å). Trans mission electron microscopy (TEM) and scanning transmission electron microscopy (STEM) investigations are undertaken at radii of 3.0 mm of the HPT deformed disks [3]. (S)TEM studies are carried out using a field emission gun transmission electron microscope (JEOL JEM-2100 F, Japan) equipped with an image-side spherical aberration corrector. The electron beam of TEM is perpendicular to the shear plane of the disks for all microstructural investigations reported in this work. Morphologies of initial specimens are examined using a Scanning Electron Microscope (SEM) LEO Gemini 1525 (Carl Zeiss, Oberkochen, Germany). Oxygen contents and element valence states in the Cu–Fe samples are
3. Results and discussion 3.1. Microstructural evolution Fig. 1a and b shows morphologies of the initial powder disks after consolidation and first-step of HPT deformation and the arc-melted bulk composite of the 75Cu–25Fe samples. Both samples exhibit a homoge nous and random distribution of Fe lamellae or particles (5 – 10 μm) embedded in the Cu matrix respectively. The 75Cu–25Fe initial powder disk possesses Fe lamellae with thickness of about 5 – 15 μm and spacing of 5 – 20 μm between Fe lamellae. Fe particles in 75Cu–25Fe arc-melted bulk composite have sizes of 5 – 10 μm with spacings of 3 – 10 μm be tween particles. The full scan XRD patterns in 2θ range of 40� – 100� of as-deformed samples with 100 rotations from both consolidated pow ders (abbreviated as powder samples) and arc-melted bulk composites (abbreviated as bulk samples) are shown in Fig. 1c. It can be seen that only peaks from a single fcc phase are present in all patterns, and each diffraction peak gradually shifts to a lower angle with an increase of Fe content for both powder and bulk samples. This indicates that Fe atoms are fully dissolved into the Cu matrix forming a single fcc phase after 100 rotations deformation (at radius of 3.0 mm, strain ε ¼ 1360). To reveal the microstructural evolution of Cu–Fe samples after HPT deformation, TEM is employed to check the microstructures at radii of 3.0 mm for all as-deformed samples with 100 rotations (strain ε ¼ 1360). Fig. 2a–d and Fig. 2e–h shows the bright-field (BF) images of Cu–Fe powder and bulk samples respectively. Clearly, the grains of all samples are nearly equiaxed. The measured average grain sizes for each sample are shown in the corresponding TEM images. It is evident that increasing Fe content significantly reduces the grain size after HPT. Especially for
Fig. 1. Morphologies of (a) 75Cu–25Fe powder sample after first-step HPT deformation and (b) raw arc-melted 75Cu–25Fe bulk composite. (c) XRD spectra of different as-deformed Cu–Fe samples with 100 rotations. 2
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Fig. 2. Bright-field images of different as-deformed Cu–Fe (a–d) powder and (e–h) bulk samples with 100 rotations.
samples with relatively less Fe content (˂ 15 at.%) the Fe refinement effect is very pronounced. Finally, grain refinement gets stabilized at certain values even with more Fe content up to 35 at.%. It can be concluded that the saturated states of Cu–Fe nanocrystalline alloys are achieved with a stable grain size of 55 – 60 nm for powder sample and 105 – 110 nm for bulk sample respectively. By comparing the grain size between powder and bulk samples, each powder sample has a much smaller grain size than the corresponding bulk sample with the same composition, which can be attributed to the oxides as will be detailed below. Besides the large difference in grain size, the morphology of the grain boundaries does not only differ between samples with different Fe content but also between powder and bulk samples for given composi tion. Increasing the Fe content obviously makes the initially sharp and regular grain boundaries blurred and curved. Furthermore, powder samples have even more blurred and curved grain boundaries compared to the bulk samples, which can be clearly observed from the TEM im ages. It is known that the increased energy associated with large volume fraction of grain boundaries results in an inherent thermal or mechanical instability (i.e. potential coarsening) of nanocrystalline alloys [15,16]. The driving force for grain growth is proportional to the surface energy of grain boundary while it is in an inverse ratio to the curvature of grain boundary [17]. In another word, alloying with Fe and introducing O can significantly improve the structural stability of Cu-based nanocrystalline alloys, which can be identified as “kinetic” and “thermodynamic” mechanisms respectively [18,19]. Kinetic stabilization refers to drag forces exerted on grain boundaries resulting from particles or solute atoms, reducing their mobility or immobilizing them entirely. One popular approach to such grain boundary “pinning” is by introducing second phase particles or precipitates, a phenomenon proposed by Smith [20] and commonly known as “Zener drag”. Thermodynamic stabiliza tion aims at reducing the energy of grain boundaries, affecting the total driving force for grain growth. Grain boundary segregation is a main approach to decreasing grain boundary energy for nanostructure stabilization. In this work, via comprehensively employing the kinetic and ther modynamic stabilization effects contributed by Fe alloying and oxides, the saturated Cu–Fe nanostructure is successfully stabilized at less than
60 nm. First, Fe atoms are fully dissolved into the Cu matrix forming either individual solute atoms or Fe-rich nanoclusters with a fcc struc ture (addressed in detail in the following sections). Both individual Fe solute atoms and Fe-rich nanoclusters could impose drag forces on grain boundaries, which can be termed as “solute drag”. On the other hand, Fe solute decoration of grain boundaries could reduce the grain boundary energy as revealed by theoretical calculations [10,21], what would refer to thermodynamic stabilization methods. Second, oxides introduced during powder processing are fragmented until breaking the chemical bonding and then dissolved into Cu matrix forming interstitials, or agglomerated at grain boundaries as oxide nanoclusters which will be revealed in the following parts. All these O interstitials inside matrix and oxide nanoclusters located at grain boundaries will facilitate the kinetic stabilization of microstructures via pinning grain boundaries [10,22]. Besides the large differences in grain sizes and grain boundary morphologies, another significant difference between powder and bulk samples is related to the twin density. Powder samples possess many visible nanotwins as shown in Fig. 2a–d by arrows compared to the bulk samples. The quantitative statistics of twin densities of all samples will be shown in the following content. The much higher twin densities in powder samples can be attributed firstly to the reduction of the stacking fault energy by dissolved oxygen as elucidated in our previous work [10], and secondly to the grain size-dependent twinning characteristic in nanocrystalline alloys [23]. Deformation via dislocation motion or twinning is quite complicated in nanocrystalline materials and usually two modes coexist depending on the grain size [24]. Different models have been tried to quantify the critical grain size where the dominant deformation mechanism will turn over [23,25–27]. Although different values are given, all these models indicate that below a certain critical grain size partial dislocations emitted from grain boundaries need a lower stress to move than lattice dislocations [27]. For the grain size range of 50 – 150 nm in our Cu–Fe materials, it can be assumed that the twinning propensity will increase with decreasing grain size according to above-mentioned models. 3.2. Fe and O effects on hardness Hardness of all as-deformed Cu–Fe powder and bulk samples with 3
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various Fe contents are shown in Fig. 3 with the change of applied strain. Pure Cu powder sample possesses a hardness of about 232 – 242 HV, while the hardness is improved to 255 – 276 HV for 95Cu–5Fe, and it finally reaches 363 – 388 HV for the composition of 65Cu–35Fe. Increasing hardness with increasing Fe content is also evident for bulk samples as shown in Fig. 3b, although for given composition the bulk sample shows a much lower hardness compared to the powder sample. Interestingly, all Cu–Fe samples display an obvious saturated softening phenomenon which means the decreasing hardness with increasing strain, although the turning point of strain is different for powder and bulk samples. Generally, the powder samples almost show a monotonic hardness decrease. By contrast, the hardness of bulk samples first shows a significant increasing and then decreasing behavior, with the turning points of strains shifting to larger values for compositions with higher Fe contents. The differences in hardness increments can be attributed to the presence of oxides in powder samples. Oxides can significantly facilitate and expedite the grain refinement of Cu and Fe grains. Hardness of 75Cu–25Fe powder samples deformed with 1 rotation and 25 rotations are shown in Fig. 3c for comparison. It can be seen that hardness increases monotonically at early deformation stage from the value of 245 HV, until reaching a saturated value around 360 HV after deformation with the strain of 80 – 150 (corresponding to 5 – 10 rota tions). The unexpected decreasing hardness with a further increasing strain as shown in Fig. 3a indicates that a more homogeneous single fcc phase structure with a random distribution of oxides particles, Fe solutes and interstitials leads to a rather softened material. As long as disloca tion plasticity is the dominant deformation mechanism, which is the case at low temperatures in single phase metallic materials, the grain refinement is limited by the competition between deformation induced grain refinement and mechanical and partly thermal assisted grain boundary migration [28–30]. In pure metals or single phase alloys this equilibrium which results in a saturation of grain refinement and hardness is obtained at strains between 10 and 30. In coarse multi-phase alloys or immiscible systems where a complete supersaturation by se vere plastic deformation can be obtained, as in the investigated system Cu–Fe, the necessary strain to obtain the steady state is usually signifi cantly larger. It depends very strongly on the initial size and distribution of the phases. In the present study such first saturation of the hardness takes place at 130 and about 600 for the powder and the bulk samples, respectively. However, in the investigated cases the hardness starts to decrease and seems to reach a saturation for the powder sample at a sheer strain of
1000, in the bulk sample such saturation is not reached even after 100 rotations. Since the softening is observed in both samples, only at somewhat different strain, the reason for this decrease is most probably a slight increase in grain size, caused by an increase in the grain boundary mobility due to the dissolution of Fe rich clusters. Unfortu nately, the effect is relatively small and a sufficient accurate measure ment of the grain size is not straight-forward to give a clear answer to this question. 3.3. Fe and O effects on grain refinement process To detect the fast grain refinement and dissolution processes at the early deformation stage, 75Cu–25Fe powder samples are deformed with relatively low strains, 13.6 (1 rotation), 136 (10 rotations) and 340 (25 rotations). Microstructural evolution can be clearly reflected from BF, high angle annular dark field (HAADF) images and electron diffraction patterns as shown in Fig. 4. It can be seen that: Firstly, grain refinement of Cu matrix proceeds extremely fast from micron-scale down to about 100 nm within 1 rotation deformation (as reflected in Fig. 4a–c). After deformation with 10 rotations (Fig. 4b), grain refinement of the matrix gets saturated with an average size of 50 – 60 nm. Secondly, Fe grains are also dramatically refined down to 50 – 100 nm after 1 rotation deformation (see regions with white borders or indicated by white ar rows in Fig. 4d). However, Fe grain size continuously becomes smaller than 30 nm after deformation of 10 rotations (Fig. 4e) because of Fe dissolution into Cu matrix. Finally, Fe is almost fully dissolved into Cu after deformation of 25 rotations except some very tiny clusters of less than 3 nm as indicated by white arrows in Fig. 4f. Fe dissolution is also demonstrated in the corresponding electron diffraction pattern in Fig. 4i, where no obvious Fe diffraction spot is observed. Thirdly, oxides have been radically fragmented down to nanoparticles of less than 10 nm after merely 1 rotation deformation as shown by blue arrows in Fig. 4d. These oxide nanoparticles are located at the grain boundaries which effectively facilitate the grain refinement of Cu matrix. This data provides a direct evidence of the kinetic stabilization effect of the oxides. Diffraction pattern in Fig. 4g also conveys the information of oxides via weak diffraction ring from CuO. Nevertheless, oxides are not detected any more in both HAADF image of Fig. 4e and electron diffraction pattern of Fig. 4h after deformation of 10 rotations. As addressed already in our previous work [10], oxides have been fragmented to extremely fine nanostructures located at grain boundaries or being completely dissolved into the matrix forming O-rich clusters or interstitials, which
Fig. 3. Hardness of as-deformed (a) powder and (b) bulk samples with 100 rotations. (c) Hardness of 75Cu–25Fe powder samples deformed with 1 and 25 rotations as a function of applied strain. 4
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Fig. 4. Microstructural evolution of 75Cu–25Fe powder samples deformed with different numbers of rotations (strains). (a–c) BF–STEM images, (d–f) HAADF–STEM images and (g–i) electron diffraction patterns of the samples deformed with (a, d, g) 1 rotation (ε ¼ 13.6), (b, e, h) 10 rotations (ε ¼ 136) and (c, f, i) 25 rotations (ε ¼ 340).
can only be detected by atom probe tomography or HRTEM (will be shown in the following part) instead of conventional TEM technique.
bulk sample in a full range of 0 – 1200 eV. Some typical Cu and Fe peaks are indexed in the profiles. Magnified spectra focusing on O 1s peaks are inset in Fig. 5a, where O 1s peaks are detected clearly for the powder samples while almost no signal can be observed for the 75Cu–25Fe bulk sample at the position of 530 – 532 eV. The XPS measurements provide a direct evidence of oxygen existence in the powder samples, while the bulk samples may only contain trace amounts which could however not be detected in the full range scan. To check the oxygen contents and element valence states carefully in as-deformed samples, fine measurements with a step of 0.05 eV are
3.4. Chemical analyses by XPS The above-mentioned microstructure and properties differences are attributed to the disparate sources, powders and arc-melted bulk com posites. XPS is employed to check the constituents and element valence states of the as-deformed samples. Fig. 5a shows the XPS spectra of asdeformed 75Cu–25Fe, 85Cu–15Fe powder samples and 75Cu–25Fe 5
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Fig. 5. (a) XPS profiles of as-deformed 75Cu–25Fe, 85Cu–15Fe powder samples, and 75Cu–25Fe bulk sample in a full range of 0 – 1200 eV. (b) Fine-scan XPS spectra of O 1s and Fe 2p3/2.
carried out on the O 1s and Fe 2p3/2 peaks as shown in Fig. 5b. It can be seen that each O 1s main peak splits into two components with one O–Fe (Cu) component at 530 eV and another at 531.5 eV. Every well resolved Fe 2p3/2 spectrum also shows multiplet splitting with second component shifting to a higher energy by 0.9 eV from the main component, and both attest the Fe atoms are substantially in zero-valent states. Herein it should be emphasized that special measures have been taken to elimi nate the possible influence from surface oxide layers as mentioned in the experimental section. Further, if the surface oxide layers still exist on the sample surfaces even after thorough ion sputtering, definitely the sat ellite peaks of Fe2þ or Fe3þ can be found at the position of 710 – 712 eV [31,32]. However, none of these signals is detected for both powder and bulk samples. Therefore, the extra peak at 531.5 eV can not be caused by a hydroxide or adsorbed oxygen on surfaces [32,33]. Biesinger et al. have done systematical XPS investigations on Ni [33] and Cr [34] based oxides, where they found extra components with 1.1 – 1.8 eV higher binding energies than the main components (530 eV). These extra components have area contributions between 20 and 40% for the whole O 1s peaks, and they are attributed to the so called “defective oxide component” [32,33] or “substitutional oxygen” [34], namely, defective sites within oxide crystal or oxygen species associated with interstitial
sites. This explanation is fully consistent with the XPS observation in our Cu–Fe nanocrystalline alloys. The main O 1s components at position of 530 eV are assigned to O–Fe(Cu) bonds because of oxides formed during powders premixing [32,35], and the extra components at 531.5 eV are from interstitial oxygen atoms dissolved inside fcc matrix by severe plastic deformation [34]. 3.5. Atomic resolution analyses by HRTEM As mentioned above, the oxide nanoparticles can not be resolved any more by conventional TEM once they have been refined to a certain scale. Although atom probe tomography can resolve atomic-scale con centration mapping, it is still insufficient in indicating accurate com positions and lattice structures of nanoparticles. HRTEM is taken to detect the nanoparticle in the as-deformed 75Cu–25Fe powder sample. Fig. 6a shows a nanoparticle of 8 – 10 nm embedded inside the sample, and it is indexed as Cu2O with an zone axis of [001] based on Fast Fourier Transform (FFT) pattern in Fig. 6b. Via the FFT pattern, inverse FFT image of a clearer Cu2O nanostructure is obtained as shown in Fig. 6c. Although grain boundaries of the matrix are not clearly visible in Fig. 6a, it can be speculated that this Cu2O nanoparticle resides at grain
Fig. 6. Atomic scale structure of a Cu2O nanoparticle in as-deformed 75Cu–25Fe powder sample with 100 rotations. (a) HRTEM image, (b) FFT pattern and (c) Inverse FFT image of the Cu2O nanoparticle formed by using FFT pattern shown in (b). 6
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boundaries because O-rich and Cu-rich nanoclusters have not been found to agglomerate together inside grains in our previous study [10]. It is known that apart from the reduced grain sizes, SPD can generate many additional nanometer-scale features inside the as-deformed sam ples, such as dislocations, nanotwins, stacking faults, nano-dimensional chemical inhomogeneity. Conventional diffraction contrast analysis for these defects using TEM doesn’t work efficiently in nanocrystalline al loys because of the small grain size. XRD line profile analysis on the other hand is an indirect method and can be influenced by many factors. Detecting defects and providing statistics of their densities based on HRTEM images is a direct and precise method for nanocrystalline alloys. Fig. 7 displays four HRTEM and HAADF–STEM images showing some typical atomic-scale structures observed in as-deformed 75Cu–25Fe powder sample. Fig. 7a shows a typical HRTEM image viewed on [001] zone axis, where many edge dislocation components can be observed. After severe plastic deformation, a huge amount of dislocations up to the order of magnitude of 1016 m 2 are generated in the materials. The measured edge dislocation densities of as-deformed Cu–Fe samples are summarized in Table 1. Edge dislocation density increases with the increasing Fe content for both powder and bulk samples. However,
powder sample almost has one order of magnitude higher of the edge dislocation density compared to the bulk sample. Fig. 7b shows a low angle (6.4� misorientation) grain boundary, and multiple nanotwins nucleated from the grain boundary and transmitted into the grains. However, twinning area necks gradually and finally terminates inside the grain, and twin steps can be observed at the twin boundaries (TBs). This indicates successive emission of partial
Table 1 Edge dislocation and twin densities of as-deformed Cu–Fe powder and bulk samples with 100 rotations calculated based on a large number of HRTEM images. Composition 95Cu–5Fe 85Cu–15Fe 75Cu–25Fe 65Cu–35Fe
Edge dislocation density ( � 1016 m 2)
Twin density ( � 107 m 1)
Powder
Bulk
Powder
Bulk
0.73 1.07 2.32 2.70
0.12 0.27 0.48 0.51
0.98 1.42 3.31 3.35
0.04 0.12 0.18 0.24
Fig. 7. HRTEM and HADDF–STEM images acquired in as-deformed 75Cu–25Fe powder sample with 100 rotations displaying typical nanofeatures in Cu–Fe nanocrystalline alloys. (a) A grain viewing from [001] zone axis with multiple dislocations. (b) A lens-shaped twinning region formed via multiple twin steps, a typical low angle grain boundary and a strain-induced phase transformed area marked with a white frame. (c) A grain viewing from [011] zone axis with multiple stacking faults (SF). (d) A HADDF–STEM image showing chemical inhomogeneity of Fe solute atoms. Fe-rich clusters are schematically illustrated using white el lipses. The insets in (a, c) are enlarged images for areas marked with white frames and the twin boundaries (TB) in (b–d) are indicated by white arrows. 7
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dislocations from grain boundaries. Those emitted later are stopped within the grain either because the external force driving the motion of the partial dislocations is not sufficiently high to drive them across the grain or because the external force is withdrawn before the partial dis locations reach the opposite grain boundary [36]. Interestingly, a region marked by a white frame displays a significantly different lattice structure compared to the rest part being on zone axis of [011]. It is locally strain-induced structure transformation, which has been also observed in Cu–Cr nanocrystalline alloys [3]. Apart, some nanotwins and stacking faults (SFs) can be observed and pass through the grain as exemplarily shown in Fig. 7c. The inset of the enlarged area marked with a frame demonstrates the mixture of TB and SFs close to the grain boundary. The statistics of twin densities of different as-deformed Cu–Fe samples are shown in Table 1 via counting lengths of TBs in a unit area. It can be seen that twin density increases with increasing Fe content, and powder samples contain much higher twin densities compared to the bulk samples with the same Fe content. It has been discussed in our previous work that although a large amount of nanotwins have been observed in Cu–Fe nanocrystalline alloys, dislocation movement is still the main deformation mechanism instead of deformation twinning [10]. Dislocation motion mechanism of deformation can also be reflected from the texture evolution as shown in electron diffraction patterns in Fig. 4g–i. Texture gets more significant when deformation is increased from 10 rotations (Fig. 4h) to 25 rotations (Fig. 4i), which is a hint of typical dislocation plasticity [28,37]. Therefore, in the grain size range of 50 – 150 nm, the nucleation of a full dislocation is easier than a partial dislocation because of a lower requirement of shear stress [23]. Decreasing grain size in this range will thus increase twinning pro pensity. On the other hand, oxygen-induced SF energy lowering effect will also facilitate the increase of twin density in powder samples [10]. Besides these atomic-scale structural defects, the chemical in homogeneity of Fe solutes is also detected via HAADF–STEM image shown in Fig. 7d. Although XRD and electron diffraction confirm that all Fe atoms have been dissolved into the Cu matrix and a single fcc phase is formed, Fe solute atoms are not distributed homogeneously but form some Fe-rich nanometer-clusters with a fcc structure as marked by el lipses in HAADF–STEM image (Fig. 7d).
5. Dislocation motion remains the main deformation mechanism for Cu–Fe nanocrystalline alloys. Both edge dislocations and twin den sities increase with the increase of Fe content. Meanwhile, powder sample possesses much higher defects densities compared to the bulk sample with the same Fe content. Our findings highlight the critical role of oxygen in nanocrystalline alloys, and combined effects of Fe alloying and intentionally incorpo rated oxygen. This could be a promising route to further manipulate microstructures, defects densities and thus mechanical properties of nanocrystalline alloys. Author contributions Jinming Guo designed and conducted the experiments, and wrote the manuscript with input from all authors. Qinqin Shao and Oliver Renk assisted in carrying out the TEM experiments and hardness measure ments, respectively. Lei Li and Yunbin He performed XPS measurements. Reinhard Pippan helped analyzing the hardness data and made helpful comments on the whole manuscript. Yunbin He and Zaoli Zhang read through the manuscript and assisted the discussion of results. Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgements Jinming Guo, Oliver Renk and Reinhard Pippan gratefully acknowledge the financial support by the European Research Council (ERC) under grant agreement No. 340185 USMS. Zaoli Zhang ac knowledges the financial support from the Austrian Science Fund (FWF) via the project No. P27034 - N20. Yunbin He acknowledges the financial support from the National Natural Science Foundation of China (Nos. 51572073, 11774082). Dr. Julian Rosalie at Erich Schmid Institute of Materials Science is thanked for taking SEM images for the initial materials.
4. Conclusion In this work, Fe alloying and oxygen addition effects in Cu–Fe nanocrystalline alloys are systematically investigated. Cu–Fe nano crystalline alloys with different Fe contents are generated via HPT from two different sources, blended powders and arc-melted bulk, which possess different oxygen contents. Microstructures, valence states and defects densities of powder and bulk samples are comparatively studied mainly via different TEM and XPS techniques. The key conclusions are as follows:
References [1] J. Guo, J. Rosalie, R. Pippan, Z. Zhang, On the phase evolution and dissolution process in Cu–Cr alloys deformed by high pressure torsion, Scr. Mater. 133 (2017) 41–44. [2] Y. Estrin, A. Vinogradov, Extreme grain refinement by severe plastic deformation: A wealth of challenging science, Acta Mater. 61 (2013) 782–817. [3] J. Guo, J.M. Rosalie, R. Pippan, Z. Zhang, Revealing the microstructural evolution in Cu–Cr nanocrystalline alloys during high pressure torsion, Mater. Sci. Eng. A 695 (2017) 350–359. [4] R.Z. Valiev, Y. Estrin, Z. Horita, T.G. Langdon, M.J. Zehetbauer, Y.T. Zhu, Fundamentals of superior properties in bulk nanoSPD materials, Mater. Res. Lett. 4 (2016) 1–21. [5] R.Z. Valiev, Y. Estrin, Z. Horita, T.G. Langdon, M.J. Zehetbauer, Y. Zhu, Producing bulk ultrafine-grained materials by severe plastic deformation: Ten years later, JOM 68 (2016) 1216–1226. [6] X. Sauvage, G. Wilde, S.V. Divinski, Z. Horita, R.Z. Valiev, Grain boundaries in ultrafine grained materials processed by severe plastic deformation and related phenomena, Mater. Sci. Eng. A 540 (2012) 1–12. [7] K.A. Darling, M. Rajagopalan, M. Komarasamy, M.A. Bhatia, B.C. Hornbuckle, R. S. Mishra, K.N. Solanki, Extreme creep resistance in a microstructurally stable nanocrystalline alloy, Nature 537 (2016) 378–381. [8] M.A. Shaik, B.R. Golla, Densification, microstructure and properties of mechanically alloyed and hot-pressed Cu–15 wt% Al alloy, J. Mater. Sci. 53 (2018) 14694–14712. [9] J. Guo, G. Haberfehlner, J. Rosalie, L. Li, M.J. Duarte, G. Kothleitner, G. Dehm, Y. He, R. Pippan, Z. Zhang, In situ atomic-scale observation of oxidation and decomposition processes in nanocrystalline alloys, Nat. Commun. 9 (2018) 946. [10] J. Guo, M.J. Duarte, Y. Zhang, A. Bachmaier, C. Gammer, G. Dehm, R. Pippan, Z. Zhang, Oxygen-mediated deformation and grain refinement in Cu–Fe nanocrystalline alloys, Acta Mater. 166 (2019) 281–293. [11] E. Ma, Alloys created between immiscible elements, Prog. Mater. Sci. 50 (2005) 413–509.
1. Oxides can reduce the achievable grain size after HPT either by nanometer-scale particles or dissolved interstitial oxygen. For example, the average grain sizes of as-deformed 75Cu–25Fe powder and bulk samples differ almost by a factor of two (58 nm and 110 nm respectively). 2. The single phase supersaturated solid solutions show a reduced hardness compared to the Cu and Fe heterostructures with similar grain sizes. The hardness of powder samples reaches a saturated value within deformation of 10 rotations, and then shows a decrease with further increasing strain. It seems that a final saturation can be obtained only at very high strains when a homogenous supersatu ration is reached. 3. Fe dissolving process in the powder samples is completed before deformation of 25 rotations (ε ¼ 340). Fe dissolution in Cu matrix is confirmed by XRD and XPS measurements, and oxygen exists either in the form of interstitial atoms or as O–Fe(Cu) complexes. 4. Cu2O nanoparticles are found at grain boundaries. 8
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[12] E. Ma, Instabilities and ductility of nanocrystalline and ultrafine-grained metals, Scr. Mater. 49 (2003) 663–668. [13] M. Dao, L. Lu, R.J. Asaro, J.T.M. De Hosson, E. Ma, Toward a quantitative understanding of mechanical behavior of nanocrystalline metals, Acta Mater. 55 (2007) 4041–4065. [14] A. Bachmaier, M. Kerber, D. Setman, R. Pippan, The formation of supersaturated solid solutions in Fe–Cu alloys deformed by high-pressure torsion, Acta Mater. 60 (2012) 860–871. [15] T. Chookajorn, H.A. Murdoch, C.A. Schuh, Design of stable nanocrystalline alloys, Science 337 (2012) 951–954. [16] M.R. He, S.K. Samudrala, G. Kim, P.J. Felfer, A.J. Breen, J.M. Cairney, D. S. Gianola, Linking stress-driven microstructural evolution in nanocrystalline aluminium with grain boundary doping of oxygen, Nat. Commun. 7 (2016), 11225. [17] D. Brandon, W.D. Kaplan, Microstructural Characterization of Materials, second ed., John Wiley & Sons Ltd., 2008. [18] J. Weissmüller, Alloy effects in nanostructures, Nanostructured Mater. 3 (1993) 261–272. [19] C.C. Koch, R.O. Scattergood, M. Saber, H. Kotan, High temperature stabilization of nanocrystalline grain size: thermodynamic versus kinetic strategies, J. Mater. Res. 28 (2013) 1785–1791. [20] C.S. Smith, Introduction to Grains, Phases, and Interfaces: An Interpretation of Microstructure, Met. Technol. (1948) 2387. [21] P.C. Millett, R.P. Selvam, A. Saxena, Stabilizing nanocrystalline materials with dopants, Acta Mater. 55 (2007) 2329–2336. [22] M. Kapoor, T. Kaub, K.A. Darling, B.L. Boyce, G.B. Thompson, An atom probe study on Nb solute partitioning and nanocrystalline grain stabilization in mechanically alloyed Cu–Nb, Acta Mater. 126 (2017) 564–575. [23] M. Chen, E. Ma, K.J. Hemker, H. Sheng, Y. Wang, X. Cheng, Deformation twinning in nanocrystalline aluminum, Science 300 (2003) 1275–1277. [24] Q. Shao, J. Guo, J. Chen, Z. Zhang, Atomic-scale investigation on the structural evolution and deformation behaviors of dual-phase Cu–Cr nanocrystalline alloys processed by high- pressure torsion, J. Alloy. Comp., 2020. Submitted for publication. [25] X.Z. Liao, J.Y. Huang, Y.T. Zhu, F. Zhou, E.J. Lavernia, Nanostructures and deformation mechanisms in a cryogenically ball-milled Al–Mg alloy, Phila. Mag. 83 (2003) 3065–3075.
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Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea
Combined Fe and O effects on microstructural evolution and strengthening in CuFe nanocrystalline alloys Jinming Guo a, b, *, Qinqin Shao a, c, Oliver Renk d, Lei Li b, Yunbin He b, Zaoli Zhang a, d, Reinhard Pippan a a
Erich Schmid Institute of Materials Science, Austrian Academy of Sciences, 8700 Leoben, Austria School of Materials Science and Engineering, Hubei University, 430062 Wuhan, China Center for High Resolution Electron Microscopy, College of Materials Science and Engineering, Hunan University, 410082 Changsha, China d Department of Materials Science, Chair of Materials Physics, Montanuniversit€ at Leoben, 8700 Leoben, Austria b c
A R T I C L E I N F O
A B S T R A C T
Keywords: Nanocrystalline alloy Severe plastic deformation High pressure torsion Transmission electron microscopy Oxygen effect Microstructural evolution
Immiscible Cu–Fe composites with various compositions are mechanically alloyed by means of high pressure torsion directly from blended powders and vacuum arc-melted bulk materials respectively, which contain different contents of oxygen. All investigated compositions are deformed to extremely large strains reaching a saturated state consisting only of a single face-centered cubic structure. It is found that Fe alloying as well as O addition reduce the achievable grain sizes. The single phase supersaturated solid solutions show a reduced hardness compared to the Cu and Fe heterostructures with similar grain sizes, reflected in a decreasing hardness with increasing strain. In addition, oxygen introduced during powder processing has a significant influence on the microstructures, defect densities and thus properties. Samples generated from powders have smaller grain sizes and much higher dislocation and twin densities than the corresponding arc-melted bulk samples with the same nominal composition. The combined effects of Fe and O on the deformation behavior and hardness of Cu–Fe nanocrystalline alloys are discussed. The current work shows that microstructure and properties can be tuned not only by alloying, but even further by incorporating oxygen, which may offer another design strategy for nanocrystalline alloys.
1. Introduction Severe Plastic Deformation (SPD) offers many advantages in pro ducing bulk nanocrystalline alloys, i. e., dissolution of second phases in usually “immiscible” systems [1], considerable grain refinement down to the nanometer range [2,3]. In addition, via specifically designing and controlling nano-scaled features induced by SPD, such as deformation twins, non-equilibrium grain boundaries, dislocation substructures, so lute segregation and clustering [4–6], many nanocrystalline materials show extraordinary properties compared to their coarse-grained coun terparts [7]. Over the past decades, Cu-based nanocrystalline alloys have drawn intensive interests because of their superior strength compared to coarse-grained pure Cu. They hold a great potential for various applications, such as welding electrodes, heat sinks and elec trical applications due to their excellent electrical and thermal con ductivity in combination with high strength, wear and corrosion resistances [8]. In our previous work, we have successfully fabricated
Cu–Cr [1,3] and Cu–Fe [9,10] nanocrystalline alloys through high pressure torsion (HPT) technique, which is one of the most convenient method of SPD and allows to apply extremely large strains to bulk ma terials. For example, we refined the coarse-grained Cu–Cr alloys down to a saturated grain size of about 15 nm by deformation up to a strain of 340. Finally, in contrast to the conventionally immiscible Cu–Cr system under equilibrium conditions, a maximum amount of 32 at.% Cu was fully dissolved into the Cr matrix forming a stable body-centered cubic structure [1,3]. However, one prominent problem resulted from the consolidation and severe deformation of powder mixtures by SPD is the unavoidable introduction of oxygen [11]. The intrusive oxygen usually can enhance hardness by stabilizing finer grain sizes, which commonly has a negative effect on ductility, causing a premature fracture of the sample even in the elastic regime [12,13]. In addition, our previous work revealed the atomic-scale formation processes of oxides during heating of oxygen-containing nanocrystalline alloys [9].
* Corresponding author. Erich Schmid Institute of Materials Science, Austrian Academy of Sciences, 8700 Leoben, Austria. E-mail addresses: [email protected], [email protected] (J. Guo). https://doi.org/10.1016/j.msea.2019.138800 Received 26 September 2019; Received in revised form 7 December 2019; Accepted 9 December 2019 Available online 10 December 2019 0921-5093/© 2019 Elsevier B.V. All rights reserved.
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Therefore, understanding the oxygen effect in severely deformed bulk nanocrystalline alloys and its synergetic mechanism with other alloying elements is crucial for the future design of high performance materials. In this work, a series of Cu–Fe alloys having different Fe contents are synthesized by HPT until a saturated state consisting only of a single face-centered cubic (fcc) phase is achieved. Two different types of initial materials, blended powders and arc-melted bulk, were inves tigated to understand the oxygen effect. Effects of Fe alloying and O addition on the resulting microstructures, valence states, defects den sities and mechanical properties are revealed via systematic comparison between the different types of samples based on massive statistics. This work poses a promising route to manipulate mechanical properties of nanocrystalline alloys by comprehensive adjustments of solute alloying and intentional oxygen incorporation.
investigated by X-ray Photoelectron Spectroscopy (XPS, ESCALAB 250Xi, Thermo Fisher Scientific, Waltham, USA), employing Al Kα ra diation with a beam size of 500 μm. All samples are fully polished with a media of ethyl alcohol and then transferred to XPS chamber immedi ately, which is followed by Ar ion sputtering with ion energy of 3 KeV for 5 min to completely remove the possible surface oxide layers. Vickers microhardness measurements are conducted on a Buehler Mircomet 5100 using a load of 300 g (HV0.3). Indents are made across the radii of the disks with a spacing of 300 μm between the indents, and average values of six individual measurements on deformed disks for each deformation condition are reported in this paper. Edge dislocation and twin densities are calculated based on statistics of at least 80 high-resolution transmission electron microscopy (HRTEM) images. Edge dislocation density is obtained via dividing the number of dissociated dislocation component 2a < 100 > in HRTEM images taken along [001] zone axes by the total area of the corre sponding grains. Similarly, twin density is calculated via dividing the total length of twin boundaries in HRTEM images taken along [011] zone axes by the total area of the corresponding grains.
2. Experimental procedures The raw materials with nominal compositions of (100–x) at.%Cu – x at.%Fe (x ¼ 0, 5, 15, 25, 35, abbreviated as (100–x)Cu–xFe) are mixed from Cu and Fe powders with purity of 99.9þ% and generated by arcmelting process using Cu and Fe ingots with purity of 99.95%. Here it should be emphasized that a two-step HPT processing technique is used to generate the powder samples, and details of this technique can be found in Ref. [14]. The initial materials for powder samples mentioned in this work refer to the 8 mm disks extracted from the large 40 mm consolidated and HPT deformed disks (1st HPT step with strain of 10 – 54) as shown in the inset of Fig. 1a. The as-generated initial materials are then HPT deformed at room temperature accompanied by air cooling during the whole process. The pressure during HPT is fixed at 7.4 GPa and the rotation speed is 0.4 rotation/min for all samples. All samples are deformed into disks with radii of 4 mm. As in HPT the applied strain is a function of the disk radius, the data shown in this work is either presented as a function of strain ε or given for a certain strain (at radius of 3.0 mm from the disk center). The applied strain ε is calculated based on equation (1) shown in our previous work [3]. For powder samples, the effective strain represents only the strain applied during second-step deformation. XRD is conducted on all samples using Smartlab X-Ray Diffractom eter (Rigaku, Japan) with Cu Kα1 radiation (λ ¼ 1.540593 Å). Trans mission electron microscopy (TEM) and scanning transmission electron microscopy (STEM) investigations are undertaken at radii of 3.0 mm of the HPT deformed disks [3]. (S)TEM studies are carried out using a field emission gun transmission electron microscope (JEOL JEM-2100 F, Japan) equipped with an image-side spherical aberration corrector. The electron beam of TEM is perpendicular to the shear plane of the disks for all microstructural investigations reported in this work. Morphologies of initial specimens are examined using a Scanning Electron Microscope (SEM) LEO Gemini 1525 (Carl Zeiss, Oberkochen, Germany). Oxygen contents and element valence states in the Cu–Fe samples are
3. Results and discussion 3.1. Microstructural evolution Fig. 1a and b shows morphologies of the initial powder disks after consolidation and first-step of HPT deformation and the arc-melted bulk composite of the 75Cu–25Fe samples. Both samples exhibit a homoge nous and random distribution of Fe lamellae or particles (5 – 10 μm) embedded in the Cu matrix respectively. The 75Cu–25Fe initial powder disk possesses Fe lamellae with thickness of about 5 – 15 μm and spacing of 5 – 20 μm between Fe lamellae. Fe particles in 75Cu–25Fe arc-melted bulk composite have sizes of 5 – 10 μm with spacings of 3 – 10 μm be tween particles. The full scan XRD patterns in 2θ range of 40� – 100� of as-deformed samples with 100 rotations from both consolidated pow ders (abbreviated as powder samples) and arc-melted bulk composites (abbreviated as bulk samples) are shown in Fig. 1c. It can be seen that only peaks from a single fcc phase are present in all patterns, and each diffraction peak gradually shifts to a lower angle with an increase of Fe content for both powder and bulk samples. This indicates that Fe atoms are fully dissolved into the Cu matrix forming a single fcc phase after 100 rotations deformation (at radius of 3.0 mm, strain ε ¼ 1360). To reveal the microstructural evolution of Cu–Fe samples after HPT deformation, TEM is employed to check the microstructures at radii of 3.0 mm for all as-deformed samples with 100 rotations (strain ε ¼ 1360). Fig. 2a–d and Fig. 2e–h shows the bright-field (BF) images of Cu–Fe powder and bulk samples respectively. Clearly, the grains of all samples are nearly equiaxed. The measured average grain sizes for each sample are shown in the corresponding TEM images. It is evident that increasing Fe content significantly reduces the grain size after HPT. Especially for
Fig. 1. Morphologies of (a) 75Cu–25Fe powder sample after first-step HPT deformation and (b) raw arc-melted 75Cu–25Fe bulk composite. (c) XRD spectra of different as-deformed Cu–Fe samples with 100 rotations. 2
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Fig. 2. Bright-field images of different as-deformed Cu–Fe (a–d) powder and (e–h) bulk samples with 100 rotations.
samples with relatively less Fe content (˂ 15 at.%) the Fe refinement effect is very pronounced. Finally, grain refinement gets stabilized at certain values even with more Fe content up to 35 at.%. It can be concluded that the saturated states of Cu–Fe nanocrystalline alloys are achieved with a stable grain size of 55 – 60 nm for powder sample and 105 – 110 nm for bulk sample respectively. By comparing the grain size between powder and bulk samples, each powder sample has a much smaller grain size than the corresponding bulk sample with the same composition, which can be attributed to the oxides as will be detailed below. Besides the large difference in grain size, the morphology of the grain boundaries does not only differ between samples with different Fe content but also between powder and bulk samples for given composi tion. Increasing the Fe content obviously makes the initially sharp and regular grain boundaries blurred and curved. Furthermore, powder samples have even more blurred and curved grain boundaries compared to the bulk samples, which can be clearly observed from the TEM im ages. It is known that the increased energy associated with large volume fraction of grain boundaries results in an inherent thermal or mechanical instability (i.e. potential coarsening) of nanocrystalline alloys [15,16]. The driving force for grain growth is proportional to the surface energy of grain boundary while it is in an inverse ratio to the curvature of grain boundary [17]. In another word, alloying with Fe and introducing O can significantly improve the structural stability of Cu-based nanocrystalline alloys, which can be identified as “kinetic” and “thermodynamic” mechanisms respectively [18,19]. Kinetic stabilization refers to drag forces exerted on grain boundaries resulting from particles or solute atoms, reducing their mobility or immobilizing them entirely. One popular approach to such grain boundary “pinning” is by introducing second phase particles or precipitates, a phenomenon proposed by Smith [20] and commonly known as “Zener drag”. Thermodynamic stabiliza tion aims at reducing the energy of grain boundaries, affecting the total driving force for grain growth. Grain boundary segregation is a main approach to decreasing grain boundary energy for nanostructure stabilization. In this work, via comprehensively employing the kinetic and ther modynamic stabilization effects contributed by Fe alloying and oxides, the saturated Cu–Fe nanostructure is successfully stabilized at less than
60 nm. First, Fe atoms are fully dissolved into the Cu matrix forming either individual solute atoms or Fe-rich nanoclusters with a fcc struc ture (addressed in detail in the following sections). Both individual Fe solute atoms and Fe-rich nanoclusters could impose drag forces on grain boundaries, which can be termed as “solute drag”. On the other hand, Fe solute decoration of grain boundaries could reduce the grain boundary energy as revealed by theoretical calculations [10,21], what would refer to thermodynamic stabilization methods. Second, oxides introduced during powder processing are fragmented until breaking the chemical bonding and then dissolved into Cu matrix forming interstitials, or agglomerated at grain boundaries as oxide nanoclusters which will be revealed in the following parts. All these O interstitials inside matrix and oxide nanoclusters located at grain boundaries will facilitate the kinetic stabilization of microstructures via pinning grain boundaries [10,22]. Besides the large differences in grain sizes and grain boundary morphologies, another significant difference between powder and bulk samples is related to the twin density. Powder samples possess many visible nanotwins as shown in Fig. 2a–d by arrows compared to the bulk samples. The quantitative statistics of twin densities of all samples will be shown in the following content. The much higher twin densities in powder samples can be attributed firstly to the reduction of the stacking fault energy by dissolved oxygen as elucidated in our previous work [10], and secondly to the grain size-dependent twinning characteristic in nanocrystalline alloys [23]. Deformation via dislocation motion or twinning is quite complicated in nanocrystalline materials and usually two modes coexist depending on the grain size [24]. Different models have been tried to quantify the critical grain size where the dominant deformation mechanism will turn over [23,25–27]. Although different values are given, all these models indicate that below a certain critical grain size partial dislocations emitted from grain boundaries need a lower stress to move than lattice dislocations [27]. For the grain size range of 50 – 150 nm in our Cu–Fe materials, it can be assumed that the twinning propensity will increase with decreasing grain size according to above-mentioned models. 3.2. Fe and O effects on hardness Hardness of all as-deformed Cu–Fe powder and bulk samples with 3
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various Fe contents are shown in Fig. 3 with the change of applied strain. Pure Cu powder sample possesses a hardness of about 232 – 242 HV, while the hardness is improved to 255 – 276 HV for 95Cu–5Fe, and it finally reaches 363 – 388 HV for the composition of 65Cu–35Fe. Increasing hardness with increasing Fe content is also evident for bulk samples as shown in Fig. 3b, although for given composition the bulk sample shows a much lower hardness compared to the powder sample. Interestingly, all Cu–Fe samples display an obvious saturated softening phenomenon which means the decreasing hardness with increasing strain, although the turning point of strain is different for powder and bulk samples. Generally, the powder samples almost show a monotonic hardness decrease. By contrast, the hardness of bulk samples first shows a significant increasing and then decreasing behavior, with the turning points of strains shifting to larger values for compositions with higher Fe contents. The differences in hardness increments can be attributed to the presence of oxides in powder samples. Oxides can significantly facilitate and expedite the grain refinement of Cu and Fe grains. Hardness of 75Cu–25Fe powder samples deformed with 1 rotation and 25 rotations are shown in Fig. 3c for comparison. It can be seen that hardness increases monotonically at early deformation stage from the value of 245 HV, until reaching a saturated value around 360 HV after deformation with the strain of 80 – 150 (corresponding to 5 – 10 rota tions). The unexpected decreasing hardness with a further increasing strain as shown in Fig. 3a indicates that a more homogeneous single fcc phase structure with a random distribution of oxides particles, Fe solutes and interstitials leads to a rather softened material. As long as disloca tion plasticity is the dominant deformation mechanism, which is the case at low temperatures in single phase metallic materials, the grain refinement is limited by the competition between deformation induced grain refinement and mechanical and partly thermal assisted grain boundary migration [28–30]. In pure metals or single phase alloys this equilibrium which results in a saturation of grain refinement and hardness is obtained at strains between 10 and 30. In coarse multi-phase alloys or immiscible systems where a complete supersaturation by se vere plastic deformation can be obtained, as in the investigated system Cu–Fe, the necessary strain to obtain the steady state is usually signifi cantly larger. It depends very strongly on the initial size and distribution of the phases. In the present study such first saturation of the hardness takes place at 130 and about 600 for the powder and the bulk samples, respectively. However, in the investigated cases the hardness starts to decrease and seems to reach a saturation for the powder sample at a sheer strain of
1000, in the bulk sample such saturation is not reached even after 100 rotations. Since the softening is observed in both samples, only at somewhat different strain, the reason for this decrease is most probably a slight increase in grain size, caused by an increase in the grain boundary mobility due to the dissolution of Fe rich clusters. Unfortu nately, the effect is relatively small and a sufficient accurate measure ment of the grain size is not straight-forward to give a clear answer to this question. 3.3. Fe and O effects on grain refinement process To detect the fast grain refinement and dissolution processes at the early deformation stage, 75Cu–25Fe powder samples are deformed with relatively low strains, 13.6 (1 rotation), 136 (10 rotations) and 340 (25 rotations). Microstructural evolution can be clearly reflected from BF, high angle annular dark field (HAADF) images and electron diffraction patterns as shown in Fig. 4. It can be seen that: Firstly, grain refinement of Cu matrix proceeds extremely fast from micron-scale down to about 100 nm within 1 rotation deformation (as reflected in Fig. 4a–c). After deformation with 10 rotations (Fig. 4b), grain refinement of the matrix gets saturated with an average size of 50 – 60 nm. Secondly, Fe grains are also dramatically refined down to 50 – 100 nm after 1 rotation deformation (see regions with white borders or indicated by white ar rows in Fig. 4d). However, Fe grain size continuously becomes smaller than 30 nm after deformation of 10 rotations (Fig. 4e) because of Fe dissolution into Cu matrix. Finally, Fe is almost fully dissolved into Cu after deformation of 25 rotations except some very tiny clusters of less than 3 nm as indicated by white arrows in Fig. 4f. Fe dissolution is also demonstrated in the corresponding electron diffraction pattern in Fig. 4i, where no obvious Fe diffraction spot is observed. Thirdly, oxides have been radically fragmented down to nanoparticles of less than 10 nm after merely 1 rotation deformation as shown by blue arrows in Fig. 4d. These oxide nanoparticles are located at the grain boundaries which effectively facilitate the grain refinement of Cu matrix. This data provides a direct evidence of the kinetic stabilization effect of the oxides. Diffraction pattern in Fig. 4g also conveys the information of oxides via weak diffraction ring from CuO. Nevertheless, oxides are not detected any more in both HAADF image of Fig. 4e and electron diffraction pattern of Fig. 4h after deformation of 10 rotations. As addressed already in our previous work [10], oxides have been fragmented to extremely fine nanostructures located at grain boundaries or being completely dissolved into the matrix forming O-rich clusters or interstitials, which
Fig. 3. Hardness of as-deformed (a) powder and (b) bulk samples with 100 rotations. (c) Hardness of 75Cu–25Fe powder samples deformed with 1 and 25 rotations as a function of applied strain. 4
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Fig. 4. Microstructural evolution of 75Cu–25Fe powder samples deformed with different numbers of rotations (strains). (a–c) BF–STEM images, (d–f) HAADF–STEM images and (g–i) electron diffraction patterns of the samples deformed with (a, d, g) 1 rotation (ε ¼ 13.6), (b, e, h) 10 rotations (ε ¼ 136) and (c, f, i) 25 rotations (ε ¼ 340).
can only be detected by atom probe tomography or HRTEM (will be shown in the following part) instead of conventional TEM technique.
bulk sample in a full range of 0 – 1200 eV. Some typical Cu and Fe peaks are indexed in the profiles. Magnified spectra focusing on O 1s peaks are inset in Fig. 5a, where O 1s peaks are detected clearly for the powder samples while almost no signal can be observed for the 75Cu–25Fe bulk sample at the position of 530 – 532 eV. The XPS measurements provide a direct evidence of oxygen existence in the powder samples, while the bulk samples may only contain trace amounts which could however not be detected in the full range scan. To check the oxygen contents and element valence states carefully in as-deformed samples, fine measurements with a step of 0.05 eV are
3.4. Chemical analyses by XPS The above-mentioned microstructure and properties differences are attributed to the disparate sources, powders and arc-melted bulk com posites. XPS is employed to check the constituents and element valence states of the as-deformed samples. Fig. 5a shows the XPS spectra of asdeformed 75Cu–25Fe, 85Cu–15Fe powder samples and 75Cu–25Fe 5
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Fig. 5. (a) XPS profiles of as-deformed 75Cu–25Fe, 85Cu–15Fe powder samples, and 75Cu–25Fe bulk sample in a full range of 0 – 1200 eV. (b) Fine-scan XPS spectra of O 1s and Fe 2p3/2.
carried out on the O 1s and Fe 2p3/2 peaks as shown in Fig. 5b. It can be seen that each O 1s main peak splits into two components with one O–Fe (Cu) component at 530 eV and another at 531.5 eV. Every well resolved Fe 2p3/2 spectrum also shows multiplet splitting with second component shifting to a higher energy by 0.9 eV from the main component, and both attest the Fe atoms are substantially in zero-valent states. Herein it should be emphasized that special measures have been taken to elimi nate the possible influence from surface oxide layers as mentioned in the experimental section. Further, if the surface oxide layers still exist on the sample surfaces even after thorough ion sputtering, definitely the sat ellite peaks of Fe2þ or Fe3þ can be found at the position of 710 – 712 eV [31,32]. However, none of these signals is detected for both powder and bulk samples. Therefore, the extra peak at 531.5 eV can not be caused by a hydroxide or adsorbed oxygen on surfaces [32,33]. Biesinger et al. have done systematical XPS investigations on Ni [33] and Cr [34] based oxides, where they found extra components with 1.1 – 1.8 eV higher binding energies than the main components (530 eV). These extra components have area contributions between 20 and 40% for the whole O 1s peaks, and they are attributed to the so called “defective oxide component” [32,33] or “substitutional oxygen” [34], namely, defective sites within oxide crystal or oxygen species associated with interstitial
sites. This explanation is fully consistent with the XPS observation in our Cu–Fe nanocrystalline alloys. The main O 1s components at position of 530 eV are assigned to O–Fe(Cu) bonds because of oxides formed during powders premixing [32,35], and the extra components at 531.5 eV are from interstitial oxygen atoms dissolved inside fcc matrix by severe plastic deformation [34]. 3.5. Atomic resolution analyses by HRTEM As mentioned above, the oxide nanoparticles can not be resolved any more by conventional TEM once they have been refined to a certain scale. Although atom probe tomography can resolve atomic-scale con centration mapping, it is still insufficient in indicating accurate com positions and lattice structures of nanoparticles. HRTEM is taken to detect the nanoparticle in the as-deformed 75Cu–25Fe powder sample. Fig. 6a shows a nanoparticle of 8 – 10 nm embedded inside the sample, and it is indexed as Cu2O with an zone axis of [001] based on Fast Fourier Transform (FFT) pattern in Fig. 6b. Via the FFT pattern, inverse FFT image of a clearer Cu2O nanostructure is obtained as shown in Fig. 6c. Although grain boundaries of the matrix are not clearly visible in Fig. 6a, it can be speculated that this Cu2O nanoparticle resides at grain
Fig. 6. Atomic scale structure of a Cu2O nanoparticle in as-deformed 75Cu–25Fe powder sample with 100 rotations. (a) HRTEM image, (b) FFT pattern and (c) Inverse FFT image of the Cu2O nanoparticle formed by using FFT pattern shown in (b). 6
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boundaries because O-rich and Cu-rich nanoclusters have not been found to agglomerate together inside grains in our previous study [10]. It is known that apart from the reduced grain sizes, SPD can generate many additional nanometer-scale features inside the as-deformed sam ples, such as dislocations, nanotwins, stacking faults, nano-dimensional chemical inhomogeneity. Conventional diffraction contrast analysis for these defects using TEM doesn’t work efficiently in nanocrystalline al loys because of the small grain size. XRD line profile analysis on the other hand is an indirect method and can be influenced by many factors. Detecting defects and providing statistics of their densities based on HRTEM images is a direct and precise method for nanocrystalline alloys. Fig. 7 displays four HRTEM and HAADF–STEM images showing some typical atomic-scale structures observed in as-deformed 75Cu–25Fe powder sample. Fig. 7a shows a typical HRTEM image viewed on [001] zone axis, where many edge dislocation components can be observed. After severe plastic deformation, a huge amount of dislocations up to the order of magnitude of 1016 m 2 are generated in the materials. The measured edge dislocation densities of as-deformed Cu–Fe samples are summarized in Table 1. Edge dislocation density increases with the increasing Fe content for both powder and bulk samples. However,
powder sample almost has one order of magnitude higher of the edge dislocation density compared to the bulk sample. Fig. 7b shows a low angle (6.4� misorientation) grain boundary, and multiple nanotwins nucleated from the grain boundary and transmitted into the grains. However, twinning area necks gradually and finally terminates inside the grain, and twin steps can be observed at the twin boundaries (TBs). This indicates successive emission of partial
Table 1 Edge dislocation and twin densities of as-deformed Cu–Fe powder and bulk samples with 100 rotations calculated based on a large number of HRTEM images. Composition 95Cu–5Fe 85Cu–15Fe 75Cu–25Fe 65Cu–35Fe
Edge dislocation density ( � 1016 m 2)
Twin density ( � 107 m 1)
Powder
Bulk
Powder
Bulk
0.73 1.07 2.32 2.70
0.12 0.27 0.48 0.51
0.98 1.42 3.31 3.35
0.04 0.12 0.18 0.24
Fig. 7. HRTEM and HADDF–STEM images acquired in as-deformed 75Cu–25Fe powder sample with 100 rotations displaying typical nanofeatures in Cu–Fe nanocrystalline alloys. (a) A grain viewing from [001] zone axis with multiple dislocations. (b) A lens-shaped twinning region formed via multiple twin steps, a typical low angle grain boundary and a strain-induced phase transformed area marked with a white frame. (c) A grain viewing from [011] zone axis with multiple stacking faults (SF). (d) A HADDF–STEM image showing chemical inhomogeneity of Fe solute atoms. Fe-rich clusters are schematically illustrated using white el lipses. The insets in (a, c) are enlarged images for areas marked with white frames and the twin boundaries (TB) in (b–d) are indicated by white arrows. 7
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dislocations from grain boundaries. Those emitted later are stopped within the grain either because the external force driving the motion of the partial dislocations is not sufficiently high to drive them across the grain or because the external force is withdrawn before the partial dis locations reach the opposite grain boundary [36]. Interestingly, a region marked by a white frame displays a significantly different lattice structure compared to the rest part being on zone axis of [011]. It is locally strain-induced structure transformation, which has been also observed in Cu–Cr nanocrystalline alloys [3]. Apart, some nanotwins and stacking faults (SFs) can be observed and pass through the grain as exemplarily shown in Fig. 7c. The inset of the enlarged area marked with a frame demonstrates the mixture of TB and SFs close to the grain boundary. The statistics of twin densities of different as-deformed Cu–Fe samples are shown in Table 1 via counting lengths of TBs in a unit area. It can be seen that twin density increases with increasing Fe content, and powder samples contain much higher twin densities compared to the bulk samples with the same Fe content. It has been discussed in our previous work that although a large amount of nanotwins have been observed in Cu–Fe nanocrystalline alloys, dislocation movement is still the main deformation mechanism instead of deformation twinning [10]. Dislocation motion mechanism of deformation can also be reflected from the texture evolution as shown in electron diffraction patterns in Fig. 4g–i. Texture gets more significant when deformation is increased from 10 rotations (Fig. 4h) to 25 rotations (Fig. 4i), which is a hint of typical dislocation plasticity [28,37]. Therefore, in the grain size range of 50 – 150 nm, the nucleation of a full dislocation is easier than a partial dislocation because of a lower requirement of shear stress [23]. Decreasing grain size in this range will thus increase twinning pro pensity. On the other hand, oxygen-induced SF energy lowering effect will also facilitate the increase of twin density in powder samples [10]. Besides these atomic-scale structural defects, the chemical in homogeneity of Fe solutes is also detected via HAADF–STEM image shown in Fig. 7d. Although XRD and electron diffraction confirm that all Fe atoms have been dissolved into the Cu matrix and a single fcc phase is formed, Fe solute atoms are not distributed homogeneously but form some Fe-rich nanometer-clusters with a fcc structure as marked by el lipses in HAADF–STEM image (Fig. 7d).
5. Dislocation motion remains the main deformation mechanism for Cu–Fe nanocrystalline alloys. Both edge dislocations and twin den sities increase with the increase of Fe content. Meanwhile, powder sample possesses much higher defects densities compared to the bulk sample with the same Fe content. Our findings highlight the critical role of oxygen in nanocrystalline alloys, and combined effects of Fe alloying and intentionally incorpo rated oxygen. This could be a promising route to further manipulate microstructures, defects densities and thus mechanical properties of nanocrystalline alloys. Author contributions Jinming Guo designed and conducted the experiments, and wrote the manuscript with input from all authors. Qinqin Shao and Oliver Renk assisted in carrying out the TEM experiments and hardness measure ments, respectively. Lei Li and Yunbin He performed XPS measurements. Reinhard Pippan helped analyzing the hardness data and made helpful comments on the whole manuscript. Yunbin He and Zaoli Zhang read through the manuscript and assisted the discussion of results. Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgements Jinming Guo, Oliver Renk and Reinhard Pippan gratefully acknowledge the financial support by the European Research Council (ERC) under grant agreement No. 340185 USMS. Zaoli Zhang ac knowledges the financial support from the Austrian Science Fund (FWF) via the project No. P27034 - N20. Yunbin He acknowledges the financial support from the National Natural Science Foundation of China (Nos. 51572073, 11774082). Dr. Julian Rosalie at Erich Schmid Institute of Materials Science is thanked for taking SEM images for the initial materials.
4. Conclusion In this work, Fe alloying and oxygen addition effects in Cu–Fe nanocrystalline alloys are systematically investigated. Cu–Fe nano crystalline alloys with different Fe contents are generated via HPT from two different sources, blended powders and arc-melted bulk, which possess different oxygen contents. Microstructures, valence states and defects densities of powder and bulk samples are comparatively studied mainly via different TEM and XPS techniques. The key conclusions are as follows:
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1. Oxides can reduce the achievable grain size after HPT either by nanometer-scale particles or dissolved interstitial oxygen. For example, the average grain sizes of as-deformed 75Cu–25Fe powder and bulk samples differ almost by a factor of two (58 nm and 110 nm respectively). 2. The single phase supersaturated solid solutions show a reduced hardness compared to the Cu and Fe heterostructures with similar grain sizes. The hardness of powder samples reaches a saturated value within deformation of 10 rotations, and then shows a decrease with further increasing strain. It seems that a final saturation can be obtained only at very high strains when a homogenous supersatu ration is reached. 3. Fe dissolving process in the powder samples is completed before deformation of 25 rotations (ε ¼ 340). Fe dissolution in Cu matrix is confirmed by XRD and XPS measurements, and oxygen exists either in the form of interstitial atoms or as O–Fe(Cu) complexes. 4. Cu2O nanoparticles are found at grain boundaries. 8
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