Materials Science & Engineering A 792 (2020) 139755

Available online 27 June 20200921-5093/© 2020 Elsevier B.V. All rights reserved.

Friction stir processing of high-entropy alloy reinforced aluminum matrix composites for mechanical properties enhancement

Junchen Li a, Yulong Li b, Feifan Wang c, Xiangchen Meng a, Long Wan a, Zhibo Dong a, Yongxian Huang a,*

a State Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology, Harbin, 150001, PR China b Beijing Institute of Space Launch Technology, Beijing, 100076, PR China c China Academy of Launch Vehicle Technology, Beijing Institute of Astronautical Systems Engineering, Beijing, 100076, PR China

A R T I C L E I N F O

Keywords: High entropy alloy Friction stir processing Composite Mechanical property Interface

A B S T R A C T

Metal matrix composites have been developed to overcome the urgent demand of light-weight design. High entropy alloys (HEAs) are newly emerged as a new kind of potential reinforcement for metals due to their outstanding physical and mechanical properties. Here, a kind of novel Al matrix composite reinforced by Al0.8CoCrFeNi HEA particles was fabricated by multi-pass friction stir processing (FSP). The incorporated HEA particles were homogeneously distributed into the composites and maintained the structural integrity. The average grain size of the FSPed composites decreased from 4.6 μm of Al matrix to 2.8 μm due to the particle- stimulated nucleation mechanism. The hardness, yield strength and ultimate tensile strength of the FSPed composites increased by 56%, 42% and 22% than that of the FSPed Al matrix with no obvious decrease of elongation. Interfacial diffusion occurred and the interfacial region is confirmed to be the Al3CoCrFeNi rather than the intermetallic phases. HEA particles have great potentials in mechanical properties enhancement for conventional light-weight alloys.

1. Introduction

Aluminum (Al) alloys, one of the lightest structural metals, have the most potentials in improving energy efficiency in aerospace, electronics, automobile and defense due to the low density, high strength-to-weight ratio, easy to recycle and outstanding properties of corrosion inhabita-tion [1]. However, conventional design and processing procedures, including alloying and thermomechanical processing are two key bot-tlenecks to furtherly broaden the industrial applications of Al alloys [2]. Aluminum matrix composites (AMCs) reinforced by ceramic particles, such as oxides, carbides, nitrides and borides, etc., have been developed to improve the strength of the Al matrix [3,4], while hard ceramic particles seriously deteriorate the ductility, toughness and machinability of AMCs [5]. The weak interfacial bonding due to the poor wettability and the radical incompatibility in physical and chemical properties be-tween ceramics and Al are detrimental to the load-transfer efficiency [6]. The ceramics-Al interfaces and the surroundings are favorable sites for the nucleation of micro cracks, thus deteriorates the fracture toughness [7]. Nanocomposites incorporated with nano-sized ceramics

in Al or Mg matrix are reported to both enhance strength and maintain ductility [8–10]. While the nano-sized reinforcements are difficult to distribute uniformly in matrix and the particle agglomeration de-teriorates the mechanical properties [11–13]. Further, the fabrication processes are complicated due to the intrinsic characteristics of nano-sized reinforcements. Hence, based on the connatural character-istics of ceramics, replacing the ceramic reinforcement and maintain the ductility and toughness of composites is extremely urgent.

High entropy alloys (HEAs), as a new concept of alloying design consisting of five or more principle elements, have been developed rapidly in the last ten years [14–16]. HEAs are fabricated in the forms of bulk structural materials [17] and functional coating or films [18,19] due to their superior properties, including ultra-high strength and ductility, good thermal stability [20] and high-temperature mechanical properties [21], impressive wear [22] and corrosion resistance [23]. Hence, HEAs have great potentials to be adopted as the reinforcement [24–26] based on the obvious advantages than ceramics reinforcements: sound interfacial bonding and excellent compatibility in physical met-allurgy with metallic matrix [27]. HEA particles reinforced light-weight

* Corresponding author. E-mail address: yxhuang@hit.edu.cn (Y. Huang).

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Materials Science & Engineering A

journal homepage: http://www.elsevier.com/locate/msea

https://doi.org/10.1016/j.msea.2020.139755 Received 11 March 2020; Received in revised form 9 June 2020; Accepted 13 June 2020

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metal matrix composites, including Al [28], Mg [29] and Cu [30] all reported to exhibit high strength, ductility and fracture toughness. For instance, Tan et al. [31] pointed out that the incorporated Al0.6CoCrFeNi particles into Al matrix amorphous composites through spark plasma sintering could significantly increase the fracture surface energy, strength and ductility. Besides that, the incorporated CrMnFeCoNi par-ticles in nickel aluminide through spark plasma sintering were reported to enhance the strength of matrix and maintain the compression strain [32]. Therefore, fabricating metallic composites reinforced with HEA particles through powder metallurgy have great potentials to broaden their engineering applications.

Different from the powder metallurgy fabrication procedures, fric-tion stir processing (FSP) is a short-route, green and energy efficient solid-state processing technique [33–35]. Under the severe thermo-mechanical effect, FSP have been successfully applied in microstructural modification of metallic materials [36–38] and fabri-cation of in-situ [39–41] and ex-situ composites [12,42,43]. Further-more, flash defects and welding thinning can be eliminated to maintain the integrity of FSPed materials [44]. To our best knowledge, metallic particles (e.g. Ti [41], Ni [45], Cu [46], Mo [47] and W [48]) reinforced AMCs fabricated by FSP have been reported while few research is re-ported about the incorporation of HEA particles into Al matrix by FSP. Meanwhile, interfacial reaction occurred and the inter-diffusion layer formed between the reinforcements and the matrix for both metallic

particles reinforced AMCs fabricated by FSP [45,46] and HEA particles reinforced composites fabricated by powder metallurgy [31,49]. Opin-ions about the effect of the inter-diffusion layer on mechanical proper-ties of composites are not consistent. FSP is known as a severe plastic deformation procedure with high strain rate, thus what about the interfacial diffusion behaviors of HEA reinforced AMCs fabricated by FSP? In the current work, AMCs reinforced by HEA particles were syn-thesized by FSP. The main objective is to characterize and evaluate the effect of the incorporated HEA particles on microstructural evolution and the mechanical properties of the FSPed AMCs. The interfacial diffusion between the incorporated HEA particles and the Al matrix was also investigated.

2. Materials and experiments

Commercial available AA5083-H111 plates (4 mm in thickness) and pre-alloyed Al–Co–Cr–Fe–Ni series HEA powders (~25 μm, d10 ¼ 6.5 μm, d50 ¼ 13.5 μm, d90 ¼ 23.5 μm, fabricated by vacuum inert gas at-omization) were utilized as the starting materials. The morphology, size distribution and the corresponding element mapping of the as-received HEA powders are illustrated in Fig. 1. The particles are approximately spherical in shape (the calculated average particle size is about 6.5 μm, as shown in Fig. 1b) and the Al, Co, Cr, Fe and Ni elements distribute uniformly without segregation. The calculated atomic ratio of the as-

Fig. 1. As-received HEA powders: (a) morphology, (b) size distribution and (c–g) the corresponding element mapping.

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received powders are listed in Table 1, in which the nominal composi-tion is Al0.8CoCrFeNi.

A row of blind holes with the diameter of 1.5 mm, depth of 3.0 mm and spacing distance of 2.0 mm were milled along the length of the workpiece. The workpiece surface was polished and cleaned with acetone before filling with HEA powders. The pre-treatment of capping was firstly carried out on the top of holes filled with powders using a pin- less tool with the shoulder of 10 mm in diameter, in order to enclose the powders into the workpiece and hinder them from splashing out during subsequent FSP. After that, the four-pass FSP was conducted along the centerline of the milled holes. Each fully overlapped FSP pass was conducted with a converse direction in order to eliminate the asymmetry of plastic strain. The rotational velocity of 800 rpm and processing speed of 50 mm/min were kept constant. The tool tilting angle with respect to Z-axis was 1.5� and the plunge depth was 0.15 mm. The FSP tool was made of H13 steel, of which the shoulder was 16 mm in diameter and the threaded cylinder pin was 6 mm in diameter and 3 mm in length. The as- received AA5083 plate was also fabricated by multi-pass FSP under the same processing parameters.

Microstructural analysis was conducted along the transverse direc-tion of FSP, including scanning electron microscopy (SEM), electron probe micro-analyzer (EPMA) and electron backscattered diffraction (EBSD). The phase components and element compositions were analyzed by X-ray diffraction (XRD) and energy dispersive spectroscopy (EDS), respectively. XRD examination mode was continuous scanning with the range of 20–90� and the scanning speed of 5�/min. The Vickers micro-hardness was measured along the centerline of the transverse direction with the load of 200 g and dwell time of 10 s. The nano-

indentation tests were conducted using a NanoTest Vantage with the load of 200 mN. Three dog-bone shaped samples with a gauge length of 25 mm, width of 4 mm and thickness of 2 mm were tested in ambient tensile test. The tensile specimens were located along the longitudinal direction of FSP and all the gauge zone were in the fine grain zone. The tensile tests were conducted with the constant speed of 1 mm/min and the fractographs were observed by SEM.

3. Results and discussion

3.1. Microstructural characterization

Fig. 2 shows the morphology of the FSPed AMCs in the stir zone (SZ). No macro defects exist and the incorporated HEA particles distribute uniformly without any agglomeration in the SZ. The severe plastic deformation during multi-pass FSP could accumulate the dispersion of the incorporated particles by inducing sufficient mixing, meanwhile each fully-overlapped FSP pass with a converse direction also eliminated the asymmetry of material flow in SZ. The calculated volume fraction of the HEA particles is 3.8% approximately. The incorporated particles maintain spherical after four pass FSP, which indicates the high thermal and mechanical stability of Al0.8CoCrFeNi HEAs.

Fig. 3a and b shows the inverse pole figure (IPF) maps of the FSPed Al alloys and the FSPed AMCs respectively. The grain morphology of the FSPed Al alloys and the FSPed AMCs are both equiaxed and dynamic recrystallization occurred due to the effect of severe plastic deformation and thermal exposure during FSP. The addition of HEA particles further promoted the grain refinement. The grain size distribution histograms of the FSPed Al alloys and the FSPed AMCs are shown in Fig. 4a and b. The grain size is ranging ~14 μm and ~7 μm for the FSPed Al alloys and the FSPed AMCs, and the average grain size are calculated to be 4.6 μm and 2.8 μm, respectively. The grain refinement of the FSPed AMCs is gov-erned by the dynamic crystallization of the Al matrix as well as the effect of the incorporated reinforcements. The incorporated particles with the

Table 1 Atomic ratio of the as-received HEA particles.

Element Al Co Cr Fe Ni

At. % 16.80 20.74 20.49 21.27 20.70

Fig. 2. Distribution of the incorporated HEA particles in the FSPed AMCs: (a) overall and (b) high magnification in the bottom of SZ.

Fig. 3. IPF maps of (a) the FSPed Al alloys and (b) the FSPed AMCs, the legend shows the IPF coloring of Al.

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size larger than 0.5 μm are considered to promote recrystallization via the particle-stimulated nucleation mechanism [50]. The reinforced particles can generate higher-density dislocations within the Al matrix due to the significant thermal mismatch, thus leading to the increase of stored energy to initiate the dynamic recrystallization process. Multi-pass FSP caused the excessive grain growth in the FSPed Al alloys

(as shown in Fig. 3a), leading to the formation of the bimodal distri-bution for the grain size diagrams in Fig. 4a. Accordingly, the incorpo-rated HEA particles have great potentials in the grain refinement of the FSPed composites.

Fig. 4. Histograms of the grain size distribution: (a) the FSPed Al alloys and (b) the FSPed AMCs.

Fig. 5. (a) Interfacial morphology and (b–f) element distribution of Al, Co, Cr, Fe and Ni in the FSPed AMCs.

Fig. 6. (a) The interfacial morphology between Al matrix and the incorporated HEA particles and (b) the corresponding EDS diagram at the interface.

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3.2. Interfacial diffusion

The interfacial diffusion behavior and interfacial structure between the reinforcements and the matrix are vitally important to the bearing capacity of the fabricated composites. Fig. 5 (a) displays the interfacial morphology between the HEA particles and the Al alloy matrix and (b-f) exhibit the corresponding EPMA results of the Al, Co, Cr, Fe and Ni el-ements respectively. The incorporated HEA particles still maintained the original spherical shape and some nano-sized granules formed after multi-pass FSP. The interfacial regions between the Al matrix and the incorporated HEA particles are continuous and compact. The excellent compatibility in physical metallurgy between the Al matrix and the HEA reinforcements together with the sufficient material flow during the multi-pass FSP both contribute to the effective interfacial bonding. It can be confirmed that the Al atoms diffused between the Al-HEA interfaces, resulting in the formation of Al transition layer with the thickness less than 1.0 μm (as shown in Fig. 5b). The diffusion of Al atoms was consistent with previous researches [18,49,51]. The Al element in the interfaces tends to segregate into the HEA particles thus leading to the formation of the Al-dilution regions, mainly due to the negative mixing enthalpy and strong binding forces of Al with other elements in Al–Co–Cr–Fe–Ni alloy systems [52].

Fig. 6 displays the detailed interfacial morphology of the FSPed

AMCs and the corresponding EDS diagram across the interfacial region (as marked by the scale bar in Fig. 6a). The interfacial layers are thin, smooth and homogeneous with the thickness ranging of ~200 nm. The incorporated HEA particles maintained the integrity without severe dilution and only inter-diffusion occurred at the interface with Al ma-trix. The atomic diffusion across the interfacial region is obvious, as shown in the concentration profiles of Fig. 6b. The diffusion fluxes of Co, Cr, Fe and Ni atoms are from HEA particles toward Al matrix, and the diffusion fluxes of Al atoms are in the inverse direction. The atomic diffusion induces the formation of interfacial zone with a certain width, in which the element concentration of Al, Co, Cr, Fe and Ni keeps almost constant (as marked by the dotted line in Fig. 6b). The EDS results of the interfacial region (as marked by the point P in Fig. 6a) illustrate the detailed chemical compositions of the incorporated reinforcement, in which the interfacial region is consisted of lower concentration of Al and Mg than matrix and lower concentration of Co, Cr, Fe and Ni than HEA particles. The compositions of Co, Cr Fe and Ni are close to equi-molar, which was consistent with the concentration profiles.

Fig. 7 exhibits the XRD patterns of the as-received HEA powders, the FSPed Al matrix and the FSPed AMCs respectively. According to the results of XRD, the original HEA particles were composed of single BCC phase and no other intermetallic phases was detected from the FSPed AMCs. The incorporated HEA particles were proved to be thermos- dynamically stable during multi-pass FSP and no occurrence of phase transformation. Combined with the results in Figs. 5–7, the interfacial region could not be the intermetallic phases composed of Al and tran-sition metals. Thus the compact interfacial bonding, the minute inter-facial diffusion and the clear interface without intermetallic phases all contribute to achieving the high interfacial load-transferring efficiency in the FSPed AMCs. According to the relevant researches about the electronic structure of HEAs [53,54], valence electron concentration (VEC) is adopted to formulate thermodynamic rules for the formation and phase stability of HEAs:

VEC¼Xn

i¼1ciðVECÞi

In which ci and ðVECÞi are the atomic fraction and the VEC of the ith elements, respectively. The BCC solid-solution phases are formed and stable when the VEC value is lower than 6.87 [55]. The VEC value of the interfacial diffusion region in Fig. 6 is calculated to be 5.81, thus the phase composition at the interfacial region is considered to be the Al3CoCrFeNi with BCC structure [56,57]. It should be noted that no interfacial products were detected in the XRD results in Fig. 7, mainly due to the overlapping of the diffraction angles for AlxCoCrFeNi (x > 1) HEA systems.

The diffusion behaviors occurred in the HEA-Al couples were

Fig. 7. XRD diagrams of the as-received HEA powders, FSPed Al alloys and FSPed AMCs.

Fig. 8. (a) Micro-hardness distribution and (b) nano-indentation plots of the FSPed Al alloys and the FSPed AMCs.

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controlled by the vacancy mechanism during multi-pass FSP [58], attributed to the concentration differences at the interface. The diffusion coefficients of Al0.8CoCrFeNi are significantly lower than their conven-tional counterparts, due to their ultra-high mixing entropy, multi-principle chemical constitution and severe lattice distortion [58, 59]. Meanwhile, the atomic fraction of Cr and Ni elements in the diffusion layers are relatively lower compared with Fe and Co, owing to the differences in sluggish diffusion effect of the component elements in Al0.8CoCrFeNi [60]. The atom radius of Al element is larger than the other transition elements in the Al–Co–Cr–Fe–Ni HEA systems, thus the addition of Al could change the original bonding of neighboring atoms in lattice. The lower atomic packing density of BCC structure at the inter-diffusion layers could accommodate the excessive addition of Al. Meanwhile the interfacial mismatch between the Al matrix with FCC structures and the incorporated HEA particles with BCC structure is relaxed attributed to the interfacial region of BCC structure with large lattice constant. The concentration gradients at the interfaces are the

driving force of solid-state diffusion, and the mechanical activation ef-fect caused by severe plastic deformation of FSP enhanced the interfacial diffusion [3]. Despite that, FSP is the solid-state plastic deformation method with relatively low heat input and the SZ is below the melting point of the Al matrix, thus the thickness of the diffusion layer in the FSPed AMCs is reported to be much thinner than that of the spark plasma sintering procedures [49]. As discussed above, the formation of BCC HEAs rather than the intermetallic phases at the interfacial regions plays an important role in bonding the Al matrix with the incorporated HEA particles.

3.3. Mechanical properties

The distribution of micro-hardness along the transverse direction of the FSPed Al alloys and the FSPed AMCs is exhibited in Fig. 8a. It is clearly that the incorporated HEA particles increased the hardness of the FSPed AMCs apparently. The average micro-hardness of the FSPed AMCs reached 125.7 HV0.2, which increased by 56.1% than that of the FSPed Al alloys (80.5 HV0.2). The nano-indentation curves of the FSPed samples are displayed in Fig. 8b. The average hardness and the Young’s modulus of the FSPed AMCs increased to 1.07 GPa and 102.4 GPa, exceeding about 37.2% and 16.2% over the FSPed Al alloys (0.78 GPa and 88.1 GPa).

The tensile flow behaviors at ambient temperature of the FSPed AMCs versus the FSPed Al alloys are exhibited in Fig. 9. The inserted table shows the values of yield strength (YS), ultimate tensile strength (UTS) and elongation (El.). It is clear that the incorporation of HEA particles enhanced the mechanical properties of the Al matrix excep-tionally. The YS and UTS of the FSPed AMCs reached 200 MPa and 371 MPa, with an increase by 42% and 22% than that of the FSPed Al alloys. While it is also worth pointing out that the El. of the FSPed AMCs decreased from 24.3% to 18.8%, which is inevitable in the high-strength metal matrix composites [61,62]. As mentioned above, the significant grain refinement, the uniformly-distributed reinforcements and the minute interfacial diffusion layers are responsible for mechanical properties enhancement. The fractographs of the FSPed samples are shown in Fig. 10. The fracture morphology of the FSPed Al alloys and the FSPed AMCs both indicated the typical ductile failure with large and

Fig. 9. Engineering stress-strain curves of the FSPed Al alloys and the FSPed AMCs.

Fig. 10. Fractographs for (a, b) the FSPed Al alloys and (c, d) the FSPed AMCs.

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deep dimples. The dimples of the FSPed Al alloys are larger and deeper, consistent with the engineering stress-strain curves in Fig. 9, indicating the reduction of ductility for the FSPed AMCs. The incorporated HEAs particles located at the bottom of the dimples without broken and detachment (as squared in Fig. 10c), indicating the compact interfacial bonding between the Al matrix and the HEAs particles. Meanwhile, the inter-diffusion layers are proved to be positive to contribute to the effective interfacial bonding.

4. Conclusions

In the present study, a novel Al matrix composite reinforced with Al0.8CoCrFeNi particles was fabricated by multi-pass FSP. Grain refine-ment was achieved and mechanical properties were enhanced due to the incorporated reinforcements. The main conclusions are drawn as follows:

(1) The incorporated HEA particles were uniformly distributed in Al matrix. The interfaces of HEA-Al were compact and the rein-forced particles maintained the structural integrity. Grain refinement was achieved via the particle-stimulated nucleation mechanism of the reinforced HEA particles.

(2) Interfacial diffusion occurred and the interfacial region is proved to be the Al3CoCrFeNi rather than the intermetallic phases, which is beneficial to the effective interfacial bonding.

(3) The yield strength, ultimate tensile strength, elongation of the FSPed composites reached 200 MPa, 371 MPa and 18.8%, with significant enhancement than that of the FSPed Al alloys. The HEA particles have great potentials in strengthening the light- weight metal matrix composites with outstanding mechanical properties.

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.

CRediT authorship contribution statement

Junchen Li: Conceptualization, Methodology, Investigation, Writing - original draft. Yulong Li: Methodology, Investigation. Feifan Wang: Methodology, Investigation. Xiangchen Meng: Methodology, Investigation. Long Wan: Writing - review & editing. Zhibo Dong: Writing - review & editing. Yongxian Huang: Conceptualization, Writing - review & editing, Supervision.

Acknowledgements

The work was supported by the National Natural Science Foundation of China (No. 51575132).

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Friction stir processing of high-entropy alloy reinforced aluminum matrix composites for mechanical properties enhancement1 Introduction2 Materials and experiments3 Results and discussion3.1 Microstructural characterization3.2 Interfacial diffusion3.3 Mechanical properties

4 ConclusionsDeclaration of competing interestCRediT authorship contribution statementAcknowledgementsReferences