Effect of B on the thermal stabilization of cryomilled nanocrystalline Cu–Al alloy

Materialia 5 (2019) 100253

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Effect of B on the thermal stabilization of cryomilled nanocrystalline Cu–Al alloy Koushik Sikdar a,∗, Avik Mahata b,∗, Somraj Chakravarty a, Mark A. Atwater c,d, Debdas Roy a,c, Carl C. Koch c a

Department of Materials and Metallurgical Engineering, National Institute of Foundry and Forge Technology, Ranchi 834003, India Department of Materials Science and Engineering, Missouri University of Science and Technology, Rolla, MO 65409, USA c Department of Materials Science and Engineering, North Carolina State University, Raleigh, NC 27606, USA d Department of Applied Engineering, Safety and Technology, Millersville University, Millersville, PA 17551, USA b

a r t i c l e

i n f o

Keywords: Nanocrystalline materials Thermal stability Mechanical property X-ray diffraction Grain growth Zener pinning

a b s t r a c t Nanocrystalline Cu86 Al12 B2 alloy with an as-milled average grain size of ∼11 nm was synthesized by high-energy ball milling at cryogenic temperature. The alloy was then annealed up to 900 °C (or 0.87 Tm of Cu). Microstructural changes with annealing were assessed by X-ray diffraction (XRD), transmission electron microscopy (TEM), and microindentation. TEM investigation indicates that the newly-developed alloy retains its nanoscale grain size after annealing at 900 °C. The present investigation complements our previous findings of the thermal stability of cryomilled binary Cu88 Al12 and Cu86 Al14 alloys, where excellent thermal stability was observed and attributed to the in-situ formation of nanoscale Cu–Al precipitates and their Zener pinning effect. First principle calculation shows equilibrium in Al–Cu–B prefers segregation of B in the grain boundary of Cu–Al alloy. Superior thermal stability of the Cu86 Al12 B2 alloy was primarily ascribed to the synergistic grain boundary pinning of CuAl2 and AlB12 intermetallic phases. Moreover, the alloy maintains a higher hardness than both Cu88 Al12 and Cu86 Al14 alloy. The grain size dependent Hall–Petch strengthening was found to be the dominant mechanism with substantial contributions from solid solution strengthening and Orowan strengthening in annealed state.

1. Introduction The microstructural refinement of metals and alloys to the nanocrystalline (nc) regime (average grain size is less than 100 nm at least in one dimension) results in significant improvements in their strength and hardness [1–3]. However, retaining these improved mechanical properties in a bulk nanostructured (solids with nanoscale or partly nanoscale microstructures) state by consolidating particulate nc-metals and alloys is difficult due to their inherent microstructural instability [3–7]. Nanoscale microstructures obtained by non-equilibrium processing techniques possess a very high-volume fraction of interfaces, triple junctions, and quadruple points [4]. The enhanced diffusivity (activation energy of grain growth is close to the grain boundary (GB) diffusion [8]) and high residual stress of nc-microstructure causes rapid grain coarsening at low homologous temperature [9,10]. The curvature dependent driving force (P) for coarsening can be quantified as [11]: 𝑃 =

∗

𝐴 𝛾0 𝑟𝐺

(1)

Where A is a constant (∼1), 𝛾 0 is the intrinsic grain boundary energy, and rG is the radius of the curvature of a grain. Moreover, the inferior thermal stability of nc materials limits their consolidation temperature using conventional techniques (e.g., hot compaction, pre-annealingcompaction-sintering, hot isostatic pressing, etc.), and, in turn, prevents their densification to the near-theoretical value. With such constraints, the as-consolidated compacts lack in sufficient inter-particle bonding. It has been reported that processing artifacts (mainly porosity) are among the main causes deteriorating the tensile ductility (normally < 2% [12]) of the bulk nc-material synthesized by a bottom-up approach. Proposed methods to thermally stabilize nc-microstructure conventionally follow two (2) avenues: kinetic mechanism and thermodynamic mechanism. In the kinetic mechanism, GB mobility is primarily restricted by particle pinning, impurity/second phase pinning, porosity drag, etc. [13–17]. On the other hand, the thermodynamic mechanism relies mainly on the reduction of total intrinsic GB energy by adding solutes with high elastic misfit, which preferentially segregate on the GBs of the parent lattice [18–21]. However, recent investigations suggest that the “micro-emulsion like” GB configuration for complete thermodynamic stabilization is rare, and the superior thermostability is believed

Corresponding authors. E-mail addresses: [email protected] (K. Sikdar), [email protected] (A. Mahata).

https://doi.org/10.1016/j.mtla.2019.100253 Received 2 October 2018; Accepted 9 February 2019 Available online 12 February 2019 2589-1529/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

K. Sikdar, A. Mahata and S. Chakravarty et al.

to be caused by the interplay and/or the hybrid contribution of both mechanisms [22–24]. The thermal stability of nc-Cu-based alloys garnered significant academic research interest due to the improvement in mechanical and physical properties without much negotiation in its electrical and thermal conductivity [25,26]. Due to its modest elastic misfit energy (ΔEel = 17.4 kJ/mol. [27]), Al is an unfavorable dopant for Cu. Although, study on the thermal stabilization of Cu by Al additions draws attention due to the presence of numerous intermetallic phases in their equilibrium phase diagram [28]. Our previous investigation on the thermostability of the Cu–Al system demonstrated that the formation of in situ nanoscale intermetallic phases in the Cu 12 at.% Al alloy pins the GBs of Cu and exhibits excellent thermal stability (hardness drops only ∼6% after annealing at 900 °C or 0.87 Tm of pure Cu) [29], but a further increase in Al concentration does not significantly contribute to an improvement [30]. The Cu–B equilibrium diagram indicates B has very low solubility in Cu (∼0.29 at.% at eutectic temperature, 1013 °C) [31]. Similarly, B has maximum solubility in Al 0.0045 at.% at 660 °C [32]. Grain boundary strengthening by Cu–B precipitates has been reported previously by Suryanarayanan et al. [33]. The addition of B to the Cu-Al alloy has drawn attention as it forms intermediate compounds with both Cu (like CuB24 and Cu3 B2 [33]) and Al (AlB2 and AlB12 [34]) during annealing. Mechanical alloying (MA) is capable of transforming the material to a high-energy (∼30 kJ/mol. [35]), nonequiilibrium state and yields a fine-grained material with extended solid solubility limit, and it was adopted for synthesizing the ternary nc Cu86 Al12 B2 alloy. Moreover, MA under cryogenic temperature in a protective (Ar) atmosphere prevents excessive cold welding, yields fine-scale material in shorter milling time with extremely low contamination. In the present work Cu86 Al12 B2 was synthesized by MA under cryogenic temperature and protective (Ar) atmosphere. The alloy was then annealed up to 900 °C to assess the stability of the nc-microstructure. Grain size and hardness values were measured as a function of annealing temperature and compared with pure nc Cu, Cu88 Al12 , and Cu86 Al14 synthesized under similar conditions. The experimentally obtained hardness value was then theoretically analyzed to provide insight on the prevailing strengthening mechanisms. In addition to this, first principle calculation was performed to gen insight into the role of B. 2. Experimental procedures and simulation method 2.1. Experimental procedure The starting materials were elemental powders of Cu, Al, and B (purity >99.9%, average particle size <10 μm). These elemental powders, in determined quantity, were loaded into 440 stainless steel milling vials (SPEX Sample Prep, Metuchen, NJ). Milling was carried out with grade 25, 440 stainless steel balls (Salem Specialty Ball) with a ball-to-powder ratio 10:1. In total, 17 balls 7.94 mm (0.3125 inch) in diameter and 16 balls of 6.35 mm (0.250 inch) in diameter were used. For comparison, nanostructured Cu was prepared by ball milling under the same conditions. The powders were stored and prepared in an argon-containing glove box (O2 < 1 ppm) and sealed within the vial before transferring to the mill. The milling process was carried out using a modified SPEX 8000 M mill/mixer under cryogenic condition (77 K) for 8 h. To maintain this condition, a specially designed nylon vial holder was used to maintain continuous flow of liquid nitrogen around the vial during milling. The major constituents (Cu and Al) are highly ductile, and cryogenic milling was preferred in order to prevent agglomeration and cold welding of the powder to the vial and milling media. Grain refinement is enhanced and achieved faster in cryomilling than room temperature milling [35]. To assess grain growth of the newly developed alloy it was subjected to isochronal annealed in a Lindberg tube furnace for 1 h at 200, 400, 600, 700, 800, and 900 °C under reducing atmosphere (5% H2 , balance Ar).

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The constituent phases of both cryomilled and annealed powders were evaluated by X-ray diffraction (XRD). Diffraction data was collected using a Rigaku DMax/A X-ray diffractometer by Cu K𝛼 radiation (𝜆 = 0.1541 nm) over the 2𝜃 range of 20–100°. In order to ensure proper instrumental alignment before each scan, the instrument was calibrated using a standard single crystal silicon sample. Diffraction patterns were smoothed by filtering (function filter) for background noise minimization and for removing K𝛼 2 peaks (optimized Rachinger’s method) by Xpowder software (http://www.xpowder.com). The crystallite size (𝑑XRD ) was calculated by applying the Scherrer equation given by [36] as: 𝑑XRD =

0.9 𝜆 𝐵 cos 𝜃𝐵

(2)

Where, B is full width half maximum (FWHM) of the most intense (111) peak, 𝜆 is the wavelength of incident X-ray radiation and 𝜃 B is Bragg’s diffraction angle. Microstructural investigation of both the as-synthesized and annealed alloy was carried out using a JEOL JEM 2000FX transmission electron microscope (TEM) at beam energy of 200 keV. TEM samples were prepared by uniaxial pressing of powders at room temperature under 2.6 GPa pressure in a 3 mm tungsten carbide die. The compacts were then thinned mechanically to 20 μm and electropolished by 30 vol.% nitric acid and methanol at −20 °C using a Tenupol-2 electropolisher. Hardness variation with annealing condition was evaluated using a Vickers microindentation tester (BuehlerMicromet II hardness tester, 25 g load and 12 s dwell time). During testing the particle diameterto-indent depth was maintained at ∼10:1. At least 10 microhardness indents were taken for each condition and averaged to get corresponding hardness values, and error bars represent one standard deviation of the collected data. 2.2. Simulation method To study the stability of B in Al–Cu matrix we also performed the first principle calculation by positioning B atoms in the Al–Cu matrix in the interstitial sites and grain boundaries. The similar studies on the effect of B on Al and Cu have been done separately by Lozovoi and Paxton [37] and Zhang et al. [38], respectively. Both studies confirmed that B not only segregates to the grain boundaries from bulk Al or Cu matrix, but its segregation also significantly enhances the grain boundary strength. Although there are no studies on segregation of B from Al–Cu matrix, B is expected to have a higher driving force of segregation due to its smaller size. However, the equilibrium energy of a set of structures with B interstitial and B at GB can be studied to investigate the differences in total energy. A 2D periodic super cell with 108 Al–Cu–B atoms is used for the calculation. 13 Al atoms is placed randomly inside the Al–Cu supercell. The energies and the lattice expansion in the direction of the GB after the relaxation have been measured. Then we create a Σ5(310)[001] grain boundary, find the segregation energies, and compare the energies with B as an interstitial and B in the GB. This procedure does not, of course, compute a guaranteed global minimum energy, but it is obviously a practical solution to a full optimization of grain boundary structure and provides support to our experimental work. The exchange-correlation contribution to the potential was included using the generalized gradient approximation (GGA) within the Perdew–Burke–Ernzerhof functional [39]. We used Quantum ESPRESSO code for our calculation. 3. Results The X-ray diffraction (XRD) patterns of Cu86 Al12 B2 powder in the as-milled and annealed (at 600, 700, 800, and 900 °C for 1 h) conditions are shown in Fig. 1. The XRD patterns consist of a single set of fcc peaks. The absence of any distinguishable peaks corresponding to elemental Al or B indicates complete dissolution of these elements into the Cu matrix after cryomilling (confirmed by TEM investigation below). In the

K. Sikdar, A. Mahata and S. Chakravarty et al.

Fig. 1. X-ray diffraction patterns of Cu86 Al12 B2 powders, 8 h cryomilled after 1 h annealing at 600, 700, 800, and 900 °C.

Fig. 2. Variation in Scherrer grain size of Cu88 Al12 , Cu86 Al14 , Cu86 Al12 B2 and pure nc-Cu up to 900 °C annealing temperature. As-milled grain size is listed as room temperature.

as-milled condition, the reduced grain size and increased lattice strain results in peak broadening and a shift of peak maxima towards lower Bragg angles as represented by the (111) peak in Fig. 1. As the annealing temperature is increased, peak intensity increases, width decreases and peak maxima position shifts toward higher Bragg angles as a result of structural relaxation and coarsening. For example, the peak maxima position of (111) peak shifts from 2𝜃 = 42.94° (as-milled) to 2𝜃 = 43.04° after annealing at 900 °C. Furthermore, the intensity of the high-angle peaks (i.e., (311) and (222)) increases noticeably with annealing temperature. Although intermetallic phases were identified in TEM (more below), XRD patterns of annealed samples do not show any trace of intermetallic phases, presumably due to its lower volume fraction. The variations in Scherrer grain size with respect to annealing temperature are shown in Fig. 2. Since this method assumes that peak broadening is due to the decrease in crystallite size only, it neglects factors such as instrumental broadening and lattice strain. Despite this, the method fits well when crystallite size is below 30–40 nm, where most data in this work was found to lie. Therefore, a grain size in the range of 20–25 nm provides a reasonably good approximation of actual value. Considering this caveat, XRD grain size of Cu86 Al12 B2 and

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Fig. 3. Variation of hardness with annealing temperature of Cu88 Al12 , Cu86 Al14 , Cu86 Al12 B2 and pure nc-Cu up to 900 °C annealing temperature. As-milled hardness is listed as room temperature.

Cu86 Al14 alloy is considered to remain nc up to 900 °C. In contrast, the grain size of pure nc Cu undergoes very rapid coarsening, and the grain size of Cu88 Al12 alloy reaches to the vicinity of accuracy limit of Scherrer method after annealing at 900 °C. Due to impurity content from milling media (stainless steel balls), pure Cu may exhibit better stability than that prepared by alternative methods (e.g., electrodeposition, inert gas condensation etc.). The variation in microhardness values with the annealing temperature are shown in Fig. 3. The hardness of as-milled Cu86 Al12 B2 (4.5 GPa) is much higher than Cu88 Al12 (2.91 GPa, ∼1.5 times higher), Cu86 Al14 (3.28 GPa, ∼1.2 times higher), and pure nc-Cu (2.88 GPa, ∼ 1.6 times higher). It might be attributed to the formation of a solid solution of B into the Cu lattice as mechanical alloying extends its solubility limit. The formation of a metastable solid solution is also supported by XRD and selected area (electron) diffraction (SAD) patterns of TEM, since no trace of any Cu-B or Al-B compound was detected. The hardness of pure nc-Cu falls rapidly with increase in annealing temperature since its hardness is attributed only to grain boundary contributions. The relative decrease in hardness of Cu86 Al12 B2 alloy at higher annealing temperature is also due to increase in grain size. The hardness of Cu88 Al12 and Cu86 Al14 after annealing 900 °C is almost equivalent to the as-milled condition. In contrast, the newly developed, B-containing alloy retains a higher hardness level throughout the annealing range. TEM investigation confirms the stability of nc structure of Cu86 Al12 B2 alloy. The TEM bright field (BF), dark field (DF) and selected area diffraction (SAD) of the as-milled Cu86 Al12 B2 alloy are shown in Fig. 4. The microstructure reveals that grain boundaries are ill-defined, with significant lattice distortion. As indicated in [40], light areas of DF-TEM are not the representative grain size as peripheral regions of grains are heavily distorted. The continuous FCC-Cu rings of SAD pattern (Fig. 4(c)) evidences that grain size is nc and consists mainly of lowangle boundaries (as no broadening of the ring can be observed). The microstructure was also observed after annealing at 700 °C. The TEM BF, DF and SAD are shown in Fig. 5. The SAD pattern (Fig. 5(c)) still shows continuous fundamental Cu rings, evidencing the stability of nc state. The unusual broadening of some places is due to the formation of high angle boundaries and fine bright spots on Cu ring represents formation of fine recrystallized grains upon annealing. Along with the fundamental Cu reflections, faint rings of intermetallic phases are also observed. Unlike previous investigations of ball milled and annealed Cu–Al alloys, CuAl2 has been found [29,30]. Additionally, the annealed

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Fig. 4. Cu86 Al12 B2 in the as-milled condition (a) bright field image, (b) dark field image and (c) corresponding SAD pattern.

alloy shows additional rings of AlB12 compound. Dispersion of secondary phases is shown in Fig. 5(d) (encircled). Duschanek and Rogl [32] reported that AlB12 is the only thermodynamically stable compound for ternary Al–M–B systems (where M = Be, Mg, V, Cr, Mn, Fe, Co, Ni, Cu). Moreover, literatures suggest precipitation of free B is not feasible, precipitation of second phase occurs as solid solution [31–33]. In order to find the equilibrium energy, a series of total energy calculations was performed for B atoms in the interstitial and GB position while gradually incrementing the Cu–Al–B unit cell size. As shown in Fig. 6(a), the Al atoms (blue) are randomly dispersed in solid solution with Cu (yellow). The B atoms (red) are placed in the GB. Each increment changes the total energy, and the starting point at can be referred to as a coherent site lattice. The position of the Al and B were kept same for all the calculations. When the size of the unit cell expanded, the energy drops (Fig. 6(b)). The minimum of the curve of the total energy of the GB system vs. the expansion corresponds to the equilibrium GB

Fig. 5. Cu86 Al12 B2 annealed for 1 h at 700 °C (a) bright field image with grain size distribution in bottom right inset, (b) dark field image (c) corresponding SAD pattern and (d) dark field image using the inner ring of SAD pattern shown in (c).

state. Without any expansion (at 0 in Fig. 6(b)) the GB remains in a coincident site lattice; however, the super cell needs to expand to find equilibrium as 13 Al atoms and 2 B atoms replace the Cu atoms. As shown in Fig. 6(b), the lattice needs much less expansion (∼0.44 Å) when B is at the GB compare to B in the interstitial site. B positioned in the interstitial size has expanded up to ∼1.10 Å for the equilibrium configuration. This confirms the conclusion made from the experimental work that, GB hollow site is larger than the interstitial sites and B atoms will be tend to move to GB from the interstitial position without causing any additional distortions to the lattice. Also, we notice significantly lower total energy for the B at GB, which indicates that there could be a strong driving force for the B atom to move from interior of the Al–Cu matrix towards the GB sites. We also notice that the total energy of system is also much less when B atoms are placed on the GB.

Fig. 6. (a) The Cu (yellow), Al (blue) and B (red) super cell with Σ5(310)[001] GB. (b) Total energies of the Cu–Al–B system with GB and interstitials with expanding unit cell size with respect to initial coincident site lattice model. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

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4. Discussion

Considering Zener–Smith criteria, k can be assumed as ∼1 for calculation of limiting grain size; however, recent studies have shown that k = 0.4 is a more suitable approximation [46]. The volume fraction of the second phase particles can be calculated from TEM images as [47]:

4.1. Microstructural stabilization First principle calculation reveals that the addition of B to Cu–Al system tends to reduce total energy upon segregation. The simulated result is in line with the conceptualization of Hondros and Seah [41,42]. Suryanarayanan et al. also claimed that, owing to the higher solubility limit of nc-GBs, B atoms tend to segregate on it [33]. In view of the above, it is therefore expected that B enhances the stability of nc-Cu–Al alloy. The improvement in microstructural stability of the present alloy has been primarily ascribed to the grain boundary pinning by nanoscale intermetallic particles (CuAl2 and/or AlB12 ). The present observation is in line with our previous investigation on thermal stability of Cu–Al alloys. In the demonstrated that thermal dissociation of as-milled solid solution gives rise to the formation of nanoscale CuAl2 and provides excellent thermal stability at high temperature (∼0.87 Tm ) [29,30]. The Zener pinning pressure (PZ ), can be estimated as [43]: 2𝑓 𝛾𝑅 𝑃𝑍 = 𝜆

(3)

Where f is the particle volume fraction, 𝛾 is the grain boundary energy, R is the degree of grain boundary-particle contact, 𝜆 is the interparticle spacing. For randomly distributed particles, R is considered as unity [44,45]. Assuming that the randomly distributed particles are spherical in shape with radius rP , Eq. (3) can be simplified as [18,46]: 𝑃𝑍 =

3𝑓 𝛾 2 𝑟𝑃

(4)

Eq. (4) indicates that an efficient grain boundary strengthening can be obtained by small particles having large volume fraction. The thermodynamic driving force (PD ) for coarsening can be given as [46]: 𝑃𝐷 =

2𝑘𝛾 𝐷avg

(5)

Where 𝐷avg is the average grain size. The kinetically stabilized limiting grain size (𝑑lim ) can be calculated by balancing Eqs. (4) and (5) as: 𝑑lim

4𝑟 =𝑘 𝑃 3𝑓

(6)

𝑓 =

0.2 𝑟𝑃 𝑑TEM

(7)

Where rP is the precipitate radius and 𝑑TEM is the grain size (obtained from TEM investigation) as 3.5 nm and 27 nm respectively (from Fig. 5). Eq. (6) yields the average grain size as 41 nm which is much (∼2 times) higher than the experimentally obtained value (27 nm) indicating that kinetically induced stabilization is aided by the secondary mechanism. 4.2. Strengthening mechanism In view of the above, the measured hardness values (H) can be attributed to grain size (Hall–Petch, 𝐻HP ), precipitates (Orowan strengthening, 𝐻Oro ) and solid solution strengthening (𝐻SS ) as: 𝐻 = 𝐻HP + 𝐻Oro + 𝐻SS

(8)

In this work, 𝐻HP was calculated by using Tabor’s relation (𝐻 = 3𝜎𝑦 [48]) as: [ ] −1∕2 𝐻HP = 3 𝜎0 + 𝑘HP 𝑑TEM (9) Where 𝜎 0 and 𝑘HP are the frictional stress and Hall–Petch slope, given as 25.5 MPa and 0.11 MN/m3/2 respectively [49] and d is the grain size (obtained from TEM micrographs of annealed samples). The 𝐻Oro value was determined as [50]: √ [ ln(2𝑟 ∕𝑟 ) ]3∕2 𝐺𝑏 ln(𝜆∕𝑟 ) 𝑃 0 0 𝐻oro = 3 3 (10) ln(𝜆∕𝑟0 ) 𝜋 𝜆 Where rP is the particle radius, r0 (= 4b) is the dislocation core radius, b is the Burgers vector (0.255 nm), G is the shear modulus (48 GPa) and 𝜆 is the inter-particle distance calculated as [51]: [√ ] 𝜋 𝜆 = 2𝑟𝑃 −1 (11) 4𝑓 In the present alloy, both Al and B are expected to contribute to solid solution strengthening. The hardening contribution of individual alloying element based on Labusch model can be estimated as ɛ4/3 c2/3

Fig. 7. (a) Contribution of B in the binary Cu–Al alloy, (b) Comparison of microhardness variation of nc-Cu and Cu–Al alloys with annealing temperature alloyed with W, Zr and Y and B. Clearly interstitial alloying maintains a higher strength level than large sized atom. See Ref. no. [55].

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(ɛ is misfit strain and c is at.% solute) [52]. However, incorporation of two different solute is difficult in this equation. In addition to this, the total amount of solute available after precipitation of Al-rich and B-rich precipitates is difficult to estimate. Therefore, the total solid solution strengthening has been determined by the method described in [24,53,54] as: 𝐻ss = 𝐻 − (𝐻HP + 𝐻Oro )

(12)

The combination of Eqs. (9), (10) and (12) results into Eq. (13), showing dependence of total hardness on grain size, Orowan and solid solution strengthening contributions as: [ ] √ [ ln(2𝑟 ∕𝑟 ) ]3∕2 𝐺𝑏 ln(𝜆∕𝑟 ) −1∕2 𝑃 0 0 𝐻 = 3 𝜎0 + 𝑘HP 𝑑TEM + 3 3 + 𝐻SS ln(𝜆∕𝑟0 ) 𝜋 𝜆 (13) The respective contributions of 𝐻HP is 2.08 GPa (∼55%), 𝐻Oro is 0.83 GPa (∼ 22%) and 𝐻SS is 0.99 GPa (∼ 23%) towards total hardness obtained from Eqs. (9), (10) and (12). It indicates that, a substantial fraction was coming from solid solution strengthening effect of B (Fig. 7(a)). Even after annealing at 900 °C, the hardness of Cu86 Al12 B2 (3.42 GPa) alloy is ∼25% higher than both Cu88 Al12 and Cu86 Al14 (∼2.73 GPa) alloy. The ratio of grain size hardening to precipitation hardening (𝐻HP : 𝐻Oro ) is ∼2.4, which shows that the thermally stabilized nc grains are the primary contributor to the strength at higher annealing temperatures. Fig. 7(b) shows the microhardness alteration of nc-Cu “alloyed” with different elements, prepared by similar technique. It is clear that the addition of B retains a higher strength level of hardness in nc Cu–Al alloy even after annealing at 900 °C. For instance, hardness of ternary Cu86 Al12 B2 alloy (3.42 GPa) is ∼22% higher than Cu86 Al12 Zr2 alloy (2.8 GPa) [27] and ∼32% higher than Cu99 Zr1 (2.6 GPa) [45] “alloy” after 900 °C annealing. The present work demonstrates a noticeable strengthening effect of nc Cu–Al alloy by only 2 at.% B addition, which can be considered useful for its high-temperature applications. 5. Conclusions The salient conclusions of present work are as follows: (1) Nanocrystalline single phase Cu86 Al12 B2 alloy was successfully synthesized by cryogenic mechanical alloying, and the nanostructured state was found to be stable even after high-temperature isochronal annealing (900 °C or 0.87 Tm of Cu). (2) A hybrid mechanism composed of kinetic pinning by nanoscale intermetallics (CuAl2 and AlB12 ) with GB energy reduction by B is considered as the reason for its superior thermal stability. (3) Addition of B significantly improves solid solution strengthening, which contributes ∼23% of the total hardness after annealing. (4) The ratio of HHP : HOro was ∼2.4 and reveals stabilized grain size contributes a major portion (∼55%) of the hardness after annealing. Acknowledgments Authors are grateful to the Indo-US Science and Technology Forum, Govt. of India, under the grant of Indo-US fellowship No IUSSTF Fellowship/2011/ 8 dated 17/3/2011 and the US Office of Naval Research under Grant no. N00014-10-1-0168. Declearation of interest statement 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.

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