Fabrication and characterization of an Al-based nanocomposite with high specific strength and good elongation using large amount CMA nanoparticles

Journal Pre-proof Fabrication and characterization of an Al-based nanocomposite with high specific strength and good elongation using large amount CMA nanoparticles H. Ramezanalizadeh PII:

S0925-8388(20)30030-X

DOI:

https://doi.org/10.1016/j.jallcom.2020.153667

Reference:

JALCOM 153667

To appear in:

Journal of Alloys and Compounds

Received Date: 28 October 2019 Revised Date:

2 January 2020

Accepted Date: 3 January 2020

Please cite this article as: H. Ramezanalizadeh, Fabrication and characterization of an Al-based nanocomposite with high specific strength and good elongation using large amount CMA nanoparticles, Journal of Alloys and Compounds (2020), doi: https://doi.org/10.1016/j.jallcom.2020.153667. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2020 Published by Elsevier B.V.

Credit Author Statement H. Ramezanalizadeh: Methodology; Writing - original draft; Writing - review & editing.

Fabrication and characterization of an Al-based nanocomposite with high specific strength and good elongation using large amount CMA nanoparticles H. Ramezanalizadeh Email: [email protected], Cell phone: +98-915-1707995 Department of Materials and Polymer Engineering, Faculty of Engineering, Hakim Sabzevari University, P.O. Box 397, Sabzevar, Iran Abstract In the present research, preparation and characterization of both powder and bulk of a new Al-CMA nanocomposite were investigated. For this, the composites were prepared by mechanical alloying/milling (MA/MM) of 10 wt. % Al3Mg2 and pure Al mix powders and then cold pressing, sintering and finally hot extrusion. According to the microscopic observations, a uniform distribution of Al3Mg2 in the Al and a clean interface between them were obtained after ball milling. In addition, analysis of particle size distribution revealed that the presence of Al3Mg2 particles and amount of process control agent had a significant influence on the composite powder refinement. In addition, the measured Carr index of 13.5% indicated that the ball milled powder offered good flowability. High-resolution microscopic studies showed that formation of twins might be one of the crystallite size decreasing mechanisms during MA/MM process. The initial powder boundary in the powder samples could not be completely deleted by cold pressing and sintering, but could be reduced by following hot extrusion. The microstructure of extruded nanocomposites comprised nanometer Al grains and Al3Mg2 nanoparticles dispersed interior grains which produced by dynamic recrystallizations (DRX). The Al3Mg2 particles experienced vigorous size reduction during MA/MM, which play an important role in Al grain refining during DRX via Zener pinning. These conditions along with clean interface for the 15 h milled sample could be resulted in notable enhancement of 515% in yield strength (YS), 603% in ultimate compressive strength (UCS), 181% in young’s modulus and 406% in microhardness compared with pure Al. A more interesting point was the elongation up to 24% owing to the uniform distribution of nanoparticles in Al grains. These results from Al-Al3Mg2 nanocomposites discover newer feasibilities for impressive manufacturing of high strength and low-density composites with improved ductility, which could make them possible candidates for a wide range of industries especially for weight critical applications.

Keywords: Aluminum matrix nanocomposites, Mechanical properties, Mechanical milling/alloying, Microstructure, Interfaces.

1. Introduction

Recently, much research work has been focused on synthesizing the composite materials because of unique combination of properties unachievable with conventional materials [1]. Amongst the metal matrix composites (MMCs), Al-based matrix composites (AMCs) attract the most attention owing to their low density, high specific strength and stiffness, excellent wear resistance etc [2]. On the other hand, metal matrix nanocomposites (MMNCs) are most promising in producing balanced mechanical properties between nano- and micro-structured materials, i.e., enhanced hardness, Young’s modulus, 0.2% yield strength, ultimate tensile strength and ductility, due to the addition of nanosized reinforcing particles into the matrix [3]. The large ductility of the Al-based matrix and the high strength of the hard reinforcements are two interesting features of AMCs. Three key factors should be considered in selection of the second phases: low density, high stiffness and good wettability between the matrix and second phase. Ceramics in the form of particles such as Al2O3, SiC, B4C, AlN, TiC, TiO2 TiB2 and CNT [4, 5], flakes or fibers, are the most common reinforcements that used in usual AMCs. Besides the high stiffness and reasonable density, ceramic materials show low wettability with the Albased matrix that leads to particles clustering, porosity and damaging interfacial reactions. Although, using a metal layers as a coat on the reinforcement surface or applying external energy during casting were proposed as solutions to improve wettability [6-8], they are not simple in practice and so poor adhesion between MMCs components is not fully resolved. Another problem in MMCs reinforced with ceramics is breaking up of ceramics, during temperature decreasing after sintering or casting, owing to the high difference in coefficient of thermal expansion (CTE) between metal matrix and ceramics. This could be the usual source of premature failure during loading. Hence, finding a novel reinforcing to replace the ceramics for AMCs is a crucial need. Complex metallic alloys (CMAs), intermetallic compounds with giant unit cells, comprising up to more than a thousand atoms per unit cell, are regarded as a new family of promising materials with high strength to weight ratio, good oxidation resistance and high temperature strength for reinforcing materials [9,10]. The β-Al3Mg2 intermetallic phase is known as a CMA with good physical and mechanical properties originated from its structure with 1168 atoms per unit cell [10]. An interesting point about this compound is its low density (2.25 g/cm3, even less than 2.7 g/cm3 for Al) in comparison to 3.95 g/cm3 for Al2O3, 3.21 g/cm3 for SiC, 2.51 g/cm3 for B4C and 3.51 g/cm3 for diamond, that normally are using as reinforcement in MMCs. Other properties

such as high hardness and wear resistance and high temperature strength along with low density make Al3Mg2 a suitable candidate for reinforcing of Al matrix composites. Powder metallurgy (PM) is recognized as an acceptable processing route to fabrication of MMCs that can minimize second phase segregation normally occurring during the casting methods. Although, during PM, agglomeration possibility of particles may not be zero yet, owing to static charges acting on powder surfaces or a high difference in size of second phase and matrix powders [11]. To solve this issue mechanical alloying/mechanical milling (MA/MM), which advanced from being a standard technique for particle size reduction in PM to important techniques for preparation of materials with improved mechanical and physical properties, has been used because of its ability in uniform distribution of particle [12,13]. This is because of deformation, fracturing and cold welding of the powder particles during MA/MM, which results in the reinforcing particles being well embedded into each matrix particle. Beside the mentioned challenges in MMCs, it is well known that the strengthening of composites is in cost of ductility, thus obtaining a balance between ductility and strength has been a long knot, especially for AMCs. According to the above statements, the purpose of present research was fabrication of a new MMNC with high performance. Therefore, it is tried to produce Al-10% wt. Al3Mg2 nanocomposite using MA/MM and sintering followed by hot extrusion processes. Since, acute perception and final control of the next consolidation process, as well as explaining the mechanical behavior obtained in the consolidated material is dependent to the powder characteristics, this study was concentrated on both powder and bulk parts. Hence, the microstructure evolution, grain refining mechanism and compressive properties of Al-10%wt. Al3Mg2 nanocomposites were discussed. An interesting phenomenon was concurrent increasing of strength and ductility in these nanocomposites. 2. Experimental procedure High commercial purity aluminum powders and Al3Mg2 nanoparticles were used as starting materials. The morphology of as-received Al and Al3Mg2 powder particles are shown in Fig. 1a and Fig. 2a, respectively. As can be seen, as-received Al powders had a random morphology, and their average particle size was calculated to be about 63 µm. The XRD pattern of as received Al3Mg2 is shown in Fig. 2b, which is representative of nanoscale reinforcement. 10 wt. % of Al3Mg2 nanoparticles were mixed with aluminum powders and mechanically milled up to 20 h to produce Al-Al3Mg2 nanocomposites. To minimize the extreme cold welding of aluminum

powders, 2 wt. % of stearic acid was used as a process control agent (PCA) in all experiments. Ball milling was executed using an attrition ball mill with rotation speed of 400 rpm and ball to powder weight ratio of 12:1. In this work, hardened chromium steel vial, containing steel balls with a diameter of 10 and 6 mm were used. The vial was evacuated and then filled with pure argon gas to prevent oxidation during the milling process. According to the milling time, the produced composites will be named, henceforth, simply as AC10-0HP, AC10-2HP, AC10-5HP, AC10-7HP, AC10-10HP, AC10-15HP and AC10-20HP for powder samples and AC10-0HB, AC10-2HB, AC10-5HB, AC10-7HB, AC10-10HB, AC10-15HB and AC10-20HB for bulk samples. In addition, P and B letters refer to powder and bulk state, respectively.

Fig. 1. Field emission SEM micrograph of the as received Al powder.

The Fe pollution of the powders during MA/MM was traced by using inductively coupled plasma (ICP) mass spectrometry. The phase compositions of samples was identified by X-ray diffractometry (XRD, Philips PW 1730, 40 kV and step of 0.02o) with Cu Kα radiation (λ = 0.15406 nm). The crystallite size of samples was estimated from the broadening of XRD peaks using Scherrer equation [14]: D = 0.9λ/BCosθ

(1)

Where D, λ, B and θ are crystallite size (nm), wave length of X-ray radiation (nm), full width at half maximum (FWHM) of main XRD peaks (rad) after removing the instrumental broadening and Bragg's diffraction angle (degree), respectively. For consolidated nanocomposites, the grain

size distribution of the Al matrix was determined from TEM images using Image-Pro Plus 6.0 software.

Fig. 2. a) The TEM picture and b) XRD graph of as-received Al3Mg2 nanoparticles. To explore the effect of reinforcing particles on the microstrucrural properties of the Al matrix, the elemental aluminum powder was also subjected to this high-energy milling (noted as A15HP sample). Particle size distribution (PSD) was determined by laser particle size analyzer (LPSA) (CILAS 1064 Liquid), to determine the volume size distribution, D10, D50 and D90 automatically. At least three samples were measured for each experiment. The apparent density of composite powders was characterized according to MPIF Standard 28, 1985. The Carr index (CI) was appliedto estimate the flow behavior of the ball-milled powders. The CI could be expressed as [15]: =

× 100%

(2)

Where VA is the apparent volume that obtains from pouring the powder into a container in the lack of any applied consolidation, and VT is the tapped volume gaining from the application of compression, for example impact or vibration. Normally, a CI of <15% is assumed to be an index of good flowability while >20% determines poor flowability [16]. The apparent and tapped volumes of the ball-milled powders were evaluated in terms of ASTM D7481-09 to measure the CI using a 100 ml standard scaled cylinder. Transmission electron microscope analysis (TEM) was performed using JEOL JEM-2010 and JEOL JEM-3010 microscopes operated at 200 kV. Field emission scanning electron microscope

(FESEM) type Carl Zeiss, Sigma, equipped with an energy dispersive spectrometer (EDS) was employed to investigate the morphology and particle size of the milled powders after different milling times. An optical microscope (OM) and field emission SEM were used to study the microstructures of extruded samples. After appropriate milling time, the milled powders were poured in a uniaxial die made of X40CrMoV51 (AISI H13). Then, as-milled powders were cold pressed at a compaction pressure of 500 MPa and then sintered at 448 ºC for 40 minutes in an accurate electrical resistance furnace with a heating rate of 10 ºC per minute. After sintering, the specimens were hot extruded at 400 º

C with extrusion ratio of 6:1. In order to avoid oxidation, the specimens were not handled from

furnace until the specimens cooled down. The densities of extruded bodies were measured by the Archimedes method according to the ASTM: B962-13. Theoretical density of compacts was calculated using the simple rule of mixtures, considering the fully dense values of aluminum and Al3Mg2. To investigate of mechanical behavior, the compression and micro indentation tests was also performed at room temperature at displacement rate of 1 mm/min. The samples for compression test were produced with 1/4 length-to-diameter ratio. Both loading surfaces of the cylindrical specimens were intently polished. Microhardness values were obtained with a Vickers indenter under a load of 100 g for 10 s. An average of 30 indentations was considered as the Vickers microhardness value. Fig. 3 represents flowchart of the experimental procedures in the present study.

Fig. 3. Flowchart of the experimental procedures.

The analysis of elastic modulus of composite was done by the ultrasonic thickness gage (Olympus, 3DL PLUS, USA) and its value was measured from the following relation [17]: E=

ρ[$(&' )) *(&+ )) ] (-' )) (-+ ))

(3)

.

Which VS is the shear velocity, VL refers to the longitudinal velocity and ρ denotes the density of composite.

3. Results and discussion 3.1. Characterization of powders during MA/MM 3.1.1. Nanocrystalline Al in nanocomposite powders For investigate the effect of milling time on the Al grain size during MA/MM, the XRD patterns of AC10-nanocomposite powders were tested, as shown in Fig. 4(a). It looks like the

existent peaks just belong to Al phase. In fact, according to the card number of 004-0787 (for Al) and 029-0048 (for Al3Mg2), from International Centre for Diffraction Data (ICDD), the main peaks of these two phases are almost located in the same 2θ. Although, it should be noted that due to minor content, small particle size and lower density of Al3Mg2 in comparison with Al, the intensity of Al3Mg2 peaks are weaker than Al ones and thus they could not be detected obviously in Al/Al3Mg2 XRD peaks. On the other hand, it is accepted that the formation of solid solution between metals causes changes in the lattice parameters. In XRD peaks, shift of peaks could be a sign of changes in the lattice parameters. In the current work, as can be observed from Fig 4(b), the main peak of Al has no significant change from its original place and thus the lattice parameters remain constant. This means that there is no formation of solid solution between Al and Al3Mg2. Meanwhile, according to Fig. 4(b), which shows the main XRD peak of samples (2θ = 38.4°), Al peaks are seen to broaden gradually due to grain refinement and lattice strain accumulation [18-21].

Fig. 4. (a) XRD patterns of AC10- nanocomposite powders ball-milled at various times and (b) the main

XRD peak of samples at 2θ = 38.4°.

Fig. 5 depicts the Al grain size resulted from the full width at half maximum (FWHM) of the main diffraction peaks in Fig. 4. With increasing milling time, the Al grain size decreases

quickly, until a fixed value about 48 nm after 15 h. This proposes that ball milling ends to a steady state after 15 h, which resulted in stable nanocrystalline Al in the nanocomposite powders. The grain decreasing to nanometer scale by MA/MM is related to the high dislocation produced by severe plastic deformation (SPD) [13]. Moreover, the addition of nanoparticles promotes the plastic deformation of the matrix and formation of nanocrystal grains [22-24] as the interaction of nanoparticles/dislocations can impel the formation of high density dislocations. Therefore, the grain size decreases rapidly at the starting until a balance between the grain refinements compel by SPD and grain coarsening by thermal recovery was attained, after which the grain size holds constant [25].

Fig. 5. Change of Al grain size with milling time in AC10-nanocomposite powders.

3.1.2. Al3Mg2 dispersion in the nanocomposite powders The most challengeable work in the nanocomposites is obtaining a homogenous dispensation of high amount nanoparticles. MA/MM is impressive in getting a uniform dispersion of nano-[2223,25-26] and micro-[24,27] particles in matrix. Fig. 6 illustrates field emission SEM pictures of Al-10 wt% Al3Mg2 powders milled for various milling times, which display two typical morphologies before and after steady state, respectively. Clearly, the nanocomposite powders are flat-shaped after 7 h milling (Fig. 6(a, b)) and equiaxed for 15 h milling times (Fig. 6(c, d)). The repeated welding-fracturing processes chip in the change from flattened to equiaxed particles, during which uniform distribution of nanoparticles was came off [26]. Further, the embedment of

large value nanoparticles in Al powders helps in increasing the Al work hardening rate and nanocomposite powder fracture trend, hence hastening the action of Al3Mg2 distribution [25]. Finally, Al3Mg2 particles are homogeneously distributed throughout the nanocomposite powders after 15 h milling, as seen in Fig. 6(e).

Fig. 6. Typical morphologies of Al-10 wt% Al3Mg2 powders after ball-milling for different time periods of (a) 7 h, (b) a magnified image of (a), (c) 15 h and (d) a maximized micrograph of (c). The (e) is an image of 15 h milled sample, which obviously shows good dispersion of embedded Al3Mg2 nanoparticles within Al particles.

The bright field TEM picture of AC10-2HP sample is shown in Fig. 7a, in which the Al3Mg2 particles have distributed all over the nanocrystalline Al matrix. In addition, Fig. 7b depicts the bright field TEM image of AC10-15HP sample, showing the more homogeneous dispersion of Al3Mg2 in the Al matrix. According to the Fig. 7b and also in comparision to initial particle size of Al3Mg2 (Fig. 2), it could be seen that the particle size of Al3Mg2 vigorously decreases after ball milling for 15 h and so could result in improvement the mechanical properties of nanocomposite after consolidation process [28]. Analysis of the some parts (point1) in Fig. 7b is given in Fig. 7c, which corresponds to the composition of Al3Mg2 nanoparticle. Also, the point analysis in other parts in the microstructure (point2 in Fig. 7b) which contain Fe in the composition (Fig. 7d), indicates the contamination of the powder by Fe during the MA/MM process. The contamination amount for the AC10-15HP nanocomposite measured about 0.2 wt. % by ICP mass spectrometry. It should be noted that, in the case of aluminum and its alloys there is a typical oxide layer that coats the powder surface [29]. Indeed, it is mentionable that the amount of these Fe and O elements is not high enough to be detected by XRD diffractometer and to affect the MMNC final products seriously.

Fig. 7. Bright field TEM micrographs of (a) AC10-2HP and (b) AC10-15HP samples, (c) chemical analysis of the point1 in (b) corresponding to the Al3Mg2 composition and (d) chemical analysis of the point2 in (b), indicating contamination of the powder by Fe during the MA/MM process.

3.1.3. Apparent density and flowability measurements Fig. 8 shows an exemplary dependence of apparent density on milling time for the Al -10 wt.% Al3Mg2 nanocomposites. 0 h of milling time refers to the apparent density of the unmilled sample. At short milling times, there is an ongoing decrease in the apparent density up to a least level at 5 h and then begins to increase with increasing milling time. At longer milling, 15 h to 20 h, the apparent density achieves a stable level, same as that of the as received powder.

The change of apparent density with milling time can be described by the variations in the morphology of the ductile particles during the process. The MA/MMed powders with low apparent densities display a laminar morphology, while those with high apparent density illustrate equiaxial morphology. Good powder packing originated from spherical morphology of the as received Al powder (Fig. 1) leads to the high initial apparent density values. The platy morphology of the MA/MMed powders at shorter time (Figs.9a) results in poorer consolidation of powder and an ongoing reduce in the apparent density amounts. In the next step of milling, the morphology of the particles shifts anew to the equiaxed type (Figs. 9b), with better packing of powder and so higher apparent density. This relation between the apparent density and the stages of MA/MM could be an index of steady state of milling process. It means that the stabilization of the apparent density indicates that the MA/MM has reached its constant state. This relationship can also be applied to assess the effect of milling and material parameters on the process.

Fig. 8. Apparent density vs. milling time for nanocomposite powders.

Fig. 9. Field emission SEM image of milled composites for: (a) 5 h, showing flattened morphology and (b) 20 h, showing equaixial type.

Composite powders' flowability is considered essential for having good mechanical properties as it affects the powder's compactability. Normally, good-flowing powders give rise to powder layers with continuous and uniform thickness (like spherical shape) while poor-flowing powders resulted in heterogeneous layers, which are harmful to the dimensional accuracy and mechanical properties of the final counterparts. The CI could be used to investigate the flow behavior of ballmilled Al-Al3Mg2 composite powders. As mentioned above, powders with a CI of <15% are considered to have good flowability while a CI >20% points weak flowability. Fig. 10 displays the flowability and CI results for pure Al, AC10-5HP and AC10-15HP composites. The apparent and tapped volumes for the pure Al were measured to be 56 mL and 48 mL, respectively. The tapped volume was obtained by tapping the 100 mL cylinder 500 times and reading the tapped volume to the closest scaled unit. Applying Eq. (2), the CI was measured to be 14.3%, 24.5% and 12.1% for pure Al, AC10-5HP and AC10-15HP, respectively.

Fig. 10. Flowability measurements and the related CI for Al, AC10-5HP and AC10-15HP samples.

The observations determined that from the three samples, the AC10-15HP composite showed the best flowability followed by the pure Al, while the AC10-5HP presented poor flowability with a CI of 25.5%. This subject ccould be described by two factors, the first of which is the powder's morphological evolution. Normally, powders with spherical shape will to have better flowability than non-spherical powders. The morphologies of the pure Al, AC10-5HP and AC1015HP composites are exhibited in Figs. 1, 9a and 6c, respectively. As seen, most of the AC105HP composite particles have a random shape while most of the pure Al and AC10-15HP powders are equiaxed in shape. The second factor is the adding of the stearic acid, which acted as a lubricant and additional enhanced the flow behavior of the MA/MMed powders by diminishing the friction between adjoining particles.

3.1.4. Effect of Al3Mg2 on the particle size of nanocomposite For studying the effect of reinforcing particles on the particle size of nanocomposite during milling process, Al powder was also milled for 15 h in the same conditions which nanocomposite was milled. Fig. 11a shows the PSD of Al powders used in this study. PSD of Al powders after 15 h of milling is shown in Fig. 11b. It is observed that mean diameter of Al powder increases with increasing milling time. This issue can be explained as follow.

It is famous that for MA/MM the presence of at least two components is crucial which one component must be harder than another one. During the milling of ductile Al powder, the cold welding process is more dominant than fracturing, so that the Al particles tend to be flattened. Therefore, the particle size of Al powder increases with increasing milling time. Fig. 11c shows the PSD of AC10-nanocomposite powder milled for 15 h. In comparison with Fig. 11b, it can be concluded that the presence of reinforcing phase could affect the particle size of nanocomposite in this study. The following explanation can describe this behavior. The cold deformation drives welding process during MA/MM. There is a critical deformation level, which lower deformation will not cause welding. The existence of hard second phase between ductile particles during MA/MM process raises local deformation near the reinforcement particles. On the other hand, reinforcing particles can be trapped in the interfacial boundaries, thus high deformation surrounds the reinforcement particle. This can be resulted in an increase in the local deformation, which improves the particle welding process. Besides, the higher local deformation inflicted by second phase increases the deformation hardening, which results in an enhancement of the fracture process. The enhancement in the welding and fracture processes owing to the existence of second phase might describe why the presence of reinforcing particles can decrease the particle size of nanocomposite [13]. Another feasible statement is that the small and hard brittle particles in the powder mixtue can act as small milling agents and so improve the system energy. As a result, a high deformation of the metallic matrix and finally reducing the particle size of nanocomposite can be obtained [26]. Furthermore, because of the pinning force of Al3Mg2 particles, the dislocation density of composite is larger than milled pure Al at similar milling time. That could speed up the decreasing of Al grain size. In addition, dislocation pinning created would prevent the grain coarsening induced by the thermal effect during MA/MM. Hence, Al grain size could be more decreased in the presence of Al3Mg2 particle.

Fig. 11. Particle size distributions of powder samples (a) as-received Al (b) A15HP, and (c) AC10-15HP.

3.1.5. Effect of PCA To attain prosperous MA/MM, it is crucial to achieve a balance between repeated cold welding and fracturing which occurred during it. In practice, this work may be hard to do for some materials (mostly soft cases) by MA/MM itself. One of these materials is Al, which a PCA, acting as a surfactant, is necessary during MA/MM [30]. Stearic acid is one of the most useful PCAs in MA/MM. Normally, for the usual MA/MM of soft metals, such as Al, 1 to 3 wt. % of stearic acid is utilized. Fig. 12a demonstrates the effect of stearic acid amount on the powder yield of MA/MM process. Powder yield here points to the weight fraction of powder collected from the milling container after completion of MA/MM. A

higher concentration of stearic acid in the milling system will ban the excessive cold welding, resulted in an improved powder yield (Fig. 12a). High powder yield for the MA/MMed composite powder is favorable due to decrease the processing cost, and so the final product cost [31]. As shown in Fig. 12a, without a PCA, only 70 wt. % of the composite powder was gathered after MA/MM, due to the severe adhesion of Al onto the milling equipments. Addition of 1 wt. % and 2 wt. % stearic acid resulted in higher powder yields of 90 wt. % and 100 wt. %, respectively. Besides improve the powder yield by stearic acid, a minimum content should be used, since it also enters some extra pollutants. Stearic acid is a carbohydrate containing oxygen, carbon and hydrogen which could introduce these elements to the powder mixes. For example, the carbon in stearic acid can produce some compounds like Al4C3 during MA/MM of AMCs [30]. These components might not be fully eliminated by the next annealing process, giving rise to the deleterious of mechanical behavior. From this point of view, stearic acid must be used as little as possible for MA/MM. Another valuable effect of PCA is on the particle size. The particle size of milled powders can be important for next densification processes. For instance, a specific particle size of milled powders is needed as a raw powder in thermal or cold spray, causing sieving of the milled powder a crucial step [32]. It should be noted that the needed PCA content during MA/MM is affected by powder composition, milling time and asked particle size [31]. In the other words, for a determined milling time and powder composition, the amount of stearic acid can directly affect the wanted particle size. Field emission SEM micrograph and PSD of AC10-15HP-1PCA sample are shown in Fig. 12b and Fig. 12c, respectively. A comparison was made between Fig. 6c and Fig. 12b and between Fig. 11c and Fig. 12c. It was seen that the use of more stearic acid could reduce the mean particle size of milled powder. The adjustment of particle size by controlling the amount of stearic acid can decrease the need for sieving.

Fig. 12. (a) Powder yield vs. stearic acid content, (b) Field emission SEM micrograph and (c) PSD of AC10-15HP-1PCA sample.

3.1.6. The interface integrity between Al and Al3Mg2 The efficiency of the densified composite samples is governed by the modality of the ingredient’s interface. It means that the failure types and mechanical properties of composites can be specified by its strength of interface. In the other words, an impressive load transfer from matrix to reinforcing could achieve with a strong interface and it results in enhanced strength and ductility. If the interface bonding is weak, interface debonding will happen when the composite is exposed to an exert bar [28].

An insight study from the interfaces between Al and Al3Mg2 particles was done by the help of TEM examinations of a selected region with details. A bright field TEM image from the AC1015HP sample is shown in Fig 13a and a SAD pattern from the determined interface by an arrow in Fig. 13a is exhibited in Fig. 13b. The SAD rings indicate a random crystal orientation of the Al and Al3Mg2 phases. The absence of any other phases in the composite confirms that the interface between matrix and second phase is clean. This image clearly indicates straight connected of Al3Mg2 lattice to the Al one without any intermediate phase. The nigh tangency of Al3Mg2 particles and Al lattice proves the lack of any unwanted phase or faults (voids or cracks) at their interface; indicating a clean interface metallurgical (Fig. 13b). The privation of reaction between Al and Al3Mg2 can be related to the all solid state during MA/MM, sintering and hot extrusion, therefore the interdiction of the interfacial reaction. These findings further prove that MA/MM, unlike other MMNC processing routs, does not lead to a feeble matrix-reinforcing interface. Worth to noting that a weak interface is harmful for mechanical efficiency. Any unwanted phases that have typically made due to inclusions were not detected in the present powder mixture. In addition, as can be seen in Fig. 13, the boundary of the Al3Mg2 particle is rather uneven, which may be created by the impact and friction during MA/MM. A jagged boundary is useful to cohesion along the interface.

Fig. 13. (a) Bright field TEM image of AC10-15HP sample and (b) SAD pattern from the determined area by an arrow in (a) showing the Al/Al3Mg2 interface is metallurgically clean.

Due to the low wettability between reinforcing surface and matrix, most of liquid state routes for preparing of AMNC need high temperatures to reach a favorable matrix-reinforcing interface. On the other hand, high temperatures produce a set of unwanted compounds that originate from the chemical reactions between matrix and second phase. These compounds are harmful for the mechanical properties. Mutually, because MA/MM is a solid state process so the presence probably of unasked compounds is low during it, nor it is in fact noticed in the EDS results (Figs. 7(c-d)) or in the HRTEM image (Fig. 13b). Therewith, the unwanted phases will not create during next densification processes, due to all of these steps accomplish at temperatures less than 450 ºC. As a result, MA/MM and hot extrusion routes able to create neat interfaces, devoid of unasked phases, between the Al matrix and the Al3Mg2 compound. MA/MM not only is capable to process of MMNCs at low temperature, to avoid the formation of unfavorable compounds, but also can leads to a powerful solid-state bond between the matrix and the second phase. This can be normally owing to the cold welding phenomenon that happens during MA/MM. This salient feature will result in strong bonds between Al and Al3Mg2 even stronger than Al/Al will, because cracks will not allow growing through the Al/Al3Mg2 interface, but instead at the matrix. As mentioned in experimental method section, a reduction ratio of 6:1 was selected for extrusion process. Moreover, reported results from the mechanical properties of the bulk MMCs produced by usual PM routes, namely mixing and extrusion without MA/MM, show the severe interface debonding during mechanical loading, denoting weak bonding between matrix and second phase [33]. Therefore, it could be concluded that the intensity of strain, which samples experienced during MA/MM, would have been eminently more than during extrusion. Thus, it is not extravagant to state that the MA/MM would have the more impact on the strength of interface than extrusion would. On the other words, cold welding, which occurs during MA/MM, could cause to a clean and strong interface between the matrix and second phase; the most important factors for load transfer effectively. 3. 1.7. Mechanism of crystallite size reduction

It is well known that the crystallite size reduction could be occurred during SPD processes such as MA/MM. In this regard, much effort has been made to study of nanostructure formation in a phenomenological way. Among the many studies, in [34], it has proposed that the creation of dislocation cells in shear bands is responsible for grain refinement during MA/MM. After this, in

[35], a dislocation model is suggested for the minimum grain size obtainable by MA/MM. In another report [36], the fatigue fracture caused by repeated strain, and cold welding is attributed to the crystallite size reduction. Fig. 14 shows the higher magnification of HRTEM image of the Al/Al3Mg2 interface. The picture inset to Fig. 14 is related selected area diffraction pattern (SADP) of rectangle region in Fig. 14. An interesting phenomenon in this figure is the presence of the twins. As can be seen from inset image, two sets of diffraction patterns could be observed and the reflection spots on the rows are split clearly. It also demonstrates a crystallographic twinning relationship with / 1) as the twinning plane. (11

Fig. 14. HRTEM image of AC10-15HP and the corresponding electron diffraction pattern (inset image).

Based on results obtained here and on information available in the literature, it could conclude that high dislocation density and formation of twins could be responsible for crystallite size refinement in this study. Revealing of this claim could be as follow descriptions. The size of initial grains is large and almost free of dislocations (Fig. 15a). With increasing structural strain caused by increasing milling time, the dislocation density raises gently (Fig. 15b). Fig. 15b illustrates less or more equiaxed grains with intense changes in contrast, representative of the faults presence such as dislocations. The dislocations could be easily observed in the inverse fast Fourier transform of HRTEM images (insets of Fig. 15b). The presence of extra half-planes in the

HRTEM images is a sign of the presence of edge dislocations after milling. By extending the MA/MM, ranging, eliminating, and rearrangement of dislocations at a given strain level will produce sub-grain boundaries. Prolonged deformation transforms the low angle sub-grain boundaries to high-angle grain boundaries. This could have done by the attraction of more dislocations into the boundaries or concomitant grain rotation during deformation. Finally, the misorientation angle of adjacent sub-grains gets to 10o, and the starting large grain converts into small grains discrete by actual boundaries. Eventually, the grain size overtakes a specific fixed level (Fig. 15c). The above statements could be schematically illustrated in Fig. 15d.

Fig. 15. Bright field TEM images of (a) AC10-2HP, (b) AC10-10HP and (c) AC10-20HP. In addition, (d) is a schematic display of grain refinement mechanism during MA/MM. The inverse fast Fourier transform images (insets of (b)) indicate edge dislocations.

On the other hand, it should be noted that the formation of nano-twins could be simplified by Al3Mg2 particles via imposing high shear strain and shear stress. This indicates that Al3Mg2 particles could act like milling medium to introduce the SPD and thus resulted in high imperfections density. Moreover, a further reduction in the crystallite size to the nanometer regime also could promote deformation twinning. That is obvious that the Al3Mg2 nano-particles can help to produce nano-crystallites in Al matrix. In other words, the presence of Al3Mg2 nano-particles could increase the performance of milling process in making nanostructure of the Al. In addition, increasing the milling time raises the dislocations density undoubtedly. This leads to further plastic deformation and imperfects density like twin boundaries. It is well known that the twin boundaries are efficient barriers for dislocation slip, thus dislocations may rapidly stack at twin boundaries and convert coherent twin boundaries into conventional grain boundaries upon major straining by MA/MM. As a result, Al3Mg2 nanoparticles and twin boundaries are two effective factors on grain size refining and therefore enhancing the mechanical properties. 3.2. Study of bulk properties

3. 2. 1. Microstructure of as-sintered body As mentioned before, MA/MM are usually utilized to produce uniformly distributed nanoparticle throughout AMCs. Although, the surfaces of milled Al powders are mostly coated by oxide films. Thus, it is crucial to break and detach these oxide films during the consolidation process [37]. The most usual techniques, which have been used to densify the milled powders into the bulk materials, are hot pressure sintering and hot extrusion. The microstructure of the as-sintered AC10-nanocomposite is illustrated in Fig. 16. The observed initial powder boundary (Fig. 16a) and a few pores plus debonding along such boundaries on the fracture surface (Fig. 16b), indicating that some oxide films and loose bonds which exist between the powder particles, is also in as-sintered nanocomposites. According to the relative density measurements of 95.2%, it can be concluded that the assintered nanocomposite are imperfectly consolidated. During conventional cold press and then

sintering, the pressure can scarcely supply high enough deformation to fully shear off the oxide films along initial powder boundary of Al [37]. Such residual powder boundary with the oxide films offends the bonding strength between the milled Al particles, therewith injuncting the material consolidation. Hence, conventional cold press and then sintering is hard to fully remove initial powder boundary and to completely densified powders.

Fig. 16. Optical (a) and field emission SEM (b) micrographs showing the microstructure of the as-sintered Al-10 wt.% Al3Mg2 nanocomposite. (a) initial powder boundaries morphology; nanoparticle clusters and pores; (b) debonding along initial powder boundaries on the fracture surface.

3. 2. 2. Microstructure of as-extruded body For more reduction of initial powder boundary and fully densification of materials, it is crucial to do the extreme deformation for the sintered nanocomposites. Fig. 17 shows the microstructures of AC10-0HB and milled AC10 nanocomposites after extrusion process. As observed in this Figure, no initial powder boundaries are found after extrusion of both unmilled and milled samples. On the other word, the first thing should be noted is occurring the many microstructural changes during extrusion. These changes include particles orientation along extrusion direction, recrystallization of matrix and porosity closure. In addition, comparison of the microstructure of samples makes it clear that uniformity in the microstructure is achievable by the use of higher milling times.

Fig. 17. Microstructures of (a) AC10-0HB, (b) AC10-2HB, (c) AC10-10HB and (d) AC10-15HB samples by optical micrograph normal to the extrusion direction.

This uniform distribution of fine particles will have a great effect on improving the mechanical properties of MMNC [28]. In fact, MA/MM through cold welding-fracture-cold welding leads to diffusion of reinforcing particles into aluminum matrix particles and prevents their separation or agglomeration in Al matrix grain boundaries [13, 26]. Fig. 18 displays the field emission SEM images of the extruded nanocomposites. These pictures show the microstructure including of dark areas (the β-Al3Mg2) homogeneously distributed in the Al matrix (the bright regions). The regions of β-Al3Mg2 become more interconnected with increasing milling time. As can be seen in Fig. 18c, only few pores are visible, further corroborating the high density of the consolidated specimens.

Fig. 18. Field emission SEM micrographs for the extruded materials (a), AC10-0HB and (b) AC10-15HB, and (c) the high magnification image from cross sectional of AC10-15HB.

It is evident that the homogeneous dispersion of the second phase is the most important thing for a composite material to show its exceptional function. The MA/MM process is one of the routes that can be applied to reach homogeneity of particle distribution throughout the matrix and reduce their size [13]. In addition, in MMCs, hot extrusion could decrease the agglomeration of second particles and so a better dispersion through the matrix. Therefore, since both MA/MM

and hot extrusion methods were used in the current investigation, the final distribution of Al3Mg2 nanoparticle would be uniform (Fig. 18c). A bright field TEM image of AC10-15HB after extrusion is depicted in Fig. 19, which further proving the uniform dispersion of the Al3Mg2 in the consolidated sample, thoroughly. By the way, relative density of 99.05% was achieved after consolidation. These findings indicate that hot extrusion can impressively reduce initial powder boundary with oxide films, the powder agglomerations and micropores. Therefore, fully consolidated composites could be achieved [13]. Furthermore, Fig. 19 shows that after MA/MM, the Al grains extended to nano-scale. During MA/MM, the powders endured ongoing plastic deformation and density of dislocations increased gently (Fig. 15b). Meantime, dislocation proliferation caused to the grain refining (Fig. 15c). It could notice that there exists a large content of fine particles with size of a few nanometers. These nanoparticles are Al3Mg2 that in principle uniformly distributed in Al matrix.

Fig. 19. Bright field TEM image of AC10-15HB sample showing an overall view of nanograins with Al3Mg2 nanopaticles distributed thorough Al matrix.

As concluded, hot extrusion could reduce micropores and consolidate the nanocomposites.

Besides, this process could help in achieving new microstructures for high strength and ductility via dynamic recrystallization (DRX). Morphology of the Al grains in the extruded AC10-15HB is shown in bright field TEM image of Fig. 20a. Recrystallized and equiaxed Al grains which

observed in Fig. 20a might be an index of occurring DRX process during hot extrusion of AC1015HB sample.

Fig 20. Bright field TEM images showing morphologies of the Al grains in the extruded AC10-15HB; (a) recrystallized Al grains, (b) Al3Mg2 nanoparticles located interior Al grains.

It is proven that DRX is a process, which in the nucleation and growth of new grains occurs during deformation by the help of driving force originated from dislocation energy. Addition of reinforcement particles may have two effects on recrystallization process: first the acceleration of it due to stimulated nucleation by particles and the second is deferring recrystallization by pinning of the grain boundaries. The shift from acceleration to delayed recrystallization normally take place when f/dp> 0.1–0.2 µm−1 [38], where f is particle volume fraction and dp is particle diameter. Here, f/dp is calculated to be 3.3 µm−1, which indicates that the addition of 10 wt. % Al3Mg2 may postpone nanocomposite DRX during hot extrusion. During grain growth, Zener pinning process may occur, which in the grain boundaries are pinned by nanoparticles. Although, the severe shear deformation during MA/MM and extrusion increases the dislocation density, which this may result in a large driving force to conquer Zenner pinning. Hence, nucleation and growth of new grains can occur even though the pinning of the grain boundaries by nanoparticle.

Fig. 20b is a high resolution of bright field TEM image that displays homogeneously dispersion of most Al3Mg2 nanoparticles within the recrystallized Al grains. As mentioned before, SPD occurred in MA/MM and hot extrusion could provide a large driving force. If this force is more than Zener pinning one is, for equalizing of that with Zener pinning force, the grain boundaries will immigrate and bypass via the nanoparticles, during grain growth. As a result, most Al3Mg2 are placed within the DRXed Al grains. Therewith, the effect of particle pinning on grain growth has been studied by 3-D phase field simulations [39]. The histogram of Al grains size distribution is presented in Fig. 21. From figure 21, it can be seen that the Al grain size is primarily dispersed in the nano scale range, with mean diameter 48 nm.

Fig. 21. The histograms of Al grain size distribution in the extruded AC10-15HB (48 nm).

The motive for nano-grains formation during DRX is high stain created by SPD processes i.e., MA/MM and hot extrusion, which can provide driving force for quantized conversion of the dislocation sub-boundaries into nano grains with large angle boundaries [40]. For high volume fractions of particles f ≥0.1, the grain diameter D can be measured based on the effect of reinforcement particles on the grain growth using Zener-pinning relation [41]: 0 = 0.728

45 67.8)

(4)

Where D is grain diameter, while dp and f are particle diameter and volume fraction, respectively. The obtained grain diameter, D =51 nm, is near to the experimentally calculated value of 48 nm, further proving that the presence of 10 wt.% Al3Mg2 nanoparticles with diameter

7 nm can truly result in the creation of nanocrystalline Al owing to the nanoparticle pinning effect on the grain growth during DRX. 3. 2. 3. Mechanical properties of hot extruded samples

Stress-elongation plots for nanocomposite and pure aluminum samples are showed in Fig. 22. The ultimate compressive strength for the nanocomposites increased with addition of the Al3Mg2 nanoparticles, and it was higher than that for pure Al. As mentioned earlier, MM produces uniform dispersion of the reinforcing particles in the matrix because repetitive fracturing and cold welding process cause the reinforcement particles to be well embedded into each aluminum particle. So, improving the compression strength by introducing Al3Mg2 nanoparticles could be explained by the homogeneous distribution of the Al3Mg2 nanoparticles thorough the aluminum matrix. The movement of dislocations could hinder by these particles in Al matrix via dispersion strengthening mechanism. Moreover, decreasing the size of Al3Mg2 nanoparticles leads to a decrease in the distance between them, which cause an increase in the required stress for dislocations movement between the Al3Mg2 nanoparticles. It consequently increases the material strength and ductility (Fig. 22) [42].

Fig. 22. (a) Compressive stress-elongation plots for pure Al and nanocomposite samples, (b) fracture surface of Al, (c) fracture surface of AC10-15HB sample.

The average yield strength (σ0.2), ultimate compressive strength (σUCS) and fracture stain (εf) are summarized in Table 1. A worth noting point was that the compressive strength and fracture strain of AC10-15HB sample were invariably higher than that of AC10-0HB one. The elastic modulus (E) of samples was also summarized in Table 1. According to the results, the milled composite possess higher elastic modulus compared with the unmilled and pure Al ones. Table 1. The elastic modulus and compressive properties of pure Al and AC10 composites. Sample E (GPa) σ0.2 (MPa) σUCS(MPa) εf (%) Al

70

80

130

24.5

AC10-0HB

72

187

300

12.1

AC10-15HB

127

412

785

24

The synchronic increment of the σUCS, ε and E can be primarily because of the enhanced interfacial bonding between Al and Al3Mg2. Amongst the conditions for improving the strength and elongation of the composites, the most important one is a good interfacial bonding which is crucial to gain effective load transfer across the Al/Al3Mg2 interface. It is well known that the E of composites could improve with the increasing of second phase amount and interfacial strength between matrix and second phase [43]. It should be reminded that, the Al3Mg2 amount was given (10% wt.), so could be not the main factor for the improved E of Al/Al3Mg2 composite. Thus, the higher E might be pointed to a stronger interfacial bonding between Al and Al3Mg2. With the higher E, higher applied force has been required for the similar elastic deformation. In addition, the applied load needed for beginning of the plastic deformation would also increase, which proves higher σ0.2 of the milled composites than that of unmilled one. The enhanced interfacial strength of milled samples compared to unmilled one, could efficiently convey the compressive load from the Al matrix to the Al3Mg2 particles during the deformation, thus the Al3Mg2 particles could be considered as a hard obstacle against the applied force. Regardless of the interfacial debonding between Al and Al3Mg2 in AC10-0HB sample, the load during the plastic deformation, which is required for multiplication, accumulation and interaction of the dislocations in AC10-15HB would be more than that of AC10-0HB. Therefore, higher compressive properties would be anticipated. Moreover, a salient interfacial strength also give rise to lower eventuality of interfacial debonding between Al and Al3Mg2, which may be the motive for the higher ductility of Al3Mg2/Al composites. Fig. 23 shows the compressive fractured surface of AC10-0HB and AC10-15HB samples. Because the interfacial decomposing could be the main source of cracks during the deformation, many debonded Al3Mg2 particles were visible on the fractured surface in Fig. 23a which is a sign of inferior ductility for unmilled composite. Therefore, it could be express that the interfacial debonding between Al and Al3Mg2 was a significant failure type of AC10-0HB composite, and the conclusion obtained here agreed with the reported works [44, 45]. However, according to the fracture study of composite milled for 15 h, Fig 23b, just a few Al3Mg2 particles were visible on the fracture surface (<10 wt.% ) and most of them are distributed within the Al matrix, revealing that the interface debonding is not primary fracture mechanism for this sample. Therefore, it could be concluded that the bonding between Al and Al3Mg2 is strong enough. Usually, a

fractography study on consolidated samples made from the MA/MMed Al-Al3Mg2 composite powder could show the interface strength indirectly. Further, fewer micropores observed on the fracture surface of AC10-15HB sample might also be one of the reasons for higher ductility of this sample than AC10-0HB. Usually, early fracture during mechanical loadings occurs due to non-uniform plasticity and premature localized deformation (necking). Naturally, if the localized deformation could be delayed, the elongation would certainly increase. It is acceptable that a higher work hardening rate is necessary for uniform plastic deformation and thus good elongation. In true stress-strain 49

49

4:

4:

curve, necking onsets when σ ≥ , which σ, ε and

refers to true stress, true strain and strain-

hardening rate, respectively [38].As one can be seen from Fig. 22, the curve slope of AC1015HB sample is obviously more, which indicates higher strain-hardening rate in comparison with AC10-0HB one. This phenomenon could be due to homogeneously dispersed nanoparticles within the Al grain (Fig. 19). A higher work hardening rate has also been reported with reducing particle size [28]. It should be mentioned that the most of particle might subject to cracking and fracturing during both milling and extrusion procedures, which could decrease particle size and interparticle spacing. This could result in high dislocation density via Orowan loops around nanoparticles and thus more strain hardening rate. Alongside the addition of Al3Mg2 nanoparticles, the proper grain size distribution of Al matrix could help in the strain hardening and therefore the elongation of Al/Al3Mg2 nanocomposites. Here, the broad grain size distribution in 20–150 nm range (see Fig. 21) and large amount of nanoparticles located interior Al grains (see Fig. 20) hinder the localized deformation and early fracture, and hence compel notable strain hardening. Finally, it is worth to note that a clean and strong interface created by MA/MM would have a great impact on the mechanical properties of the bulk composite, because of a better load transfer from matrix to the second phase [28].

Fig 23. Field emission SEM images of the compressive fracture surface for (a) AC10-0HB and (b) AC1015HB samples.

Fig 24 shows the typical microindentation effect of AC10-15HB sample. As for this sample (195 HV), the measured average micro-hardness increased by 406% and 267% in comparison with the pure Al and unmilled composite samples, respectively. This can be explained by the sample's microstructural changes. Firstly, the observed full width at half maximum of the 15 h milled sample was larger than that of the other samples (Fig. 4), which indicated the average grain size of the 15 h milled was smaller than the other samples (Fig. 5). The TEM (Figs. 7b and 19) and field emission SEM (Fig. 18c) images indicated that the Al3Mg2 reinforcement was distributed uniformly amongst the Al matrix up to 15 h of milling. Further, the distance between Al3Mg2 particulates decreased with the continued milling process. Therefore, both the effect of the ball milling on the Al matrix and the effect of the Al3Mg2 reinforcement contributed to the increase in the hardness of composite samples.

Fig. 24. Typical microindentation effect of AC10-15HB.

4. Conclusions

Al-10 wt.% Al3Mg2 nanocomposites were successfully prepared by the MA/MM, cold press, sintering and hot extrusion processes, respectively. The evolution of the microstructural, physical and mechanical features of nanocomposites was analysed using XRD, field emission SEM, HRTEM, LPSA, compression and microindentation tests. The some primary conclusions can be expressed as follows: 1. 15 h of milling was recognized as the time of steady state for MA/MM of Al-10 wt.% Al3Mg2 nanocomposite. The Al3Mg2 particles experienced a large reduction in size, and were uniformly embedded in the nanocrystalline Al matrix; minimum segregation or agglomeration was observed at this time.

2. A clean and strong interface between the Al3Mg2 particle and the Al matrix was formed during MA/MM process; no voids or cracks were observed around the interfaces. 3. Hot extrusion not only improved the density of nanocomposites, but also resulted in nano Al matrix grains via dynamic recrystallization (DRX), slight porosity and good interfacial bonding between Al and Al3Mg2. 4. Uniform dispersion of Al3Mg2 nanoparticles and refinement of Al matrix grains had considerable effect on the improvement of ultimate compressive strength (up to 603%), , yield strength (up to 515%), young’s modulus (up to 181%) and hardness (up to 406%) of the extruded

composites, while saving a good ductility up to 24%. The Al3Mg2 nanoparticles had been explored as a strong strengthening parameter of the Al matrix. References

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• • • • •

Mechanical alloying/milling and hot extrusion techniques were used to synthesize advanced aluminum-based nanocomposites with large amount of Al3Mg2 nanoparticles. A uniform distribution of the nano-reinforcements in the matrix was successfully obtained. Twins formation during milling may be one of crystallite size decreasing mechanisms. 15h milled Al-10%wtAl3Mg2 nanocomposites exhibited superior hardness, specific compressive strength, ductility and elastic modulus. Hot extrusion produced nano Al matrix grains via dynamic recrystallization (DRX). Furthermore, extruded nanocomposites exhibited minimal porosity and good interfacial bonding between Al/Al3Mg2.

Declaration of interests 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. ☐The authors declare the following financial interests/personal relationships which may be considered as potential competing interests: