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SiC composites via friction stir processing
Journal of Alloys and Compounds 798 (2019) 82e92
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Development and characterization of Al5083-CNTs/SiC composites via friction stir processing Vikram Kumar S. Jain a, K.U. Yazar b, S. Muthukumaran a, * a b
Department of Metallurgical and Materials Engineering, National Institute of Technology, Tiruchirappalli, Tamil Nadu, 620 015, India Department of Materials Engineering, Indian Institute of Science, Bengaluru, 560012, India
a r t i c l e i n f o
a b s t r a c t
Article history: Received 7 March 2019 Received in revised form 26 April 2019 Accepted 19 May 2019 Available online 21 May 2019
Carbon nanotubes (CNTs) and or micron-sized silicon carbide (SiC) particles were reinforced with Al5083 matrix to develop mono and hybrid composites via friction stir processing (FSP). The effect of CNTs/SiC either individually or in aggregate form, on microstructural evolution, texture, and mechanical properties of friction stir processed (FSPed) Al5083 composites were studied. EBSD and TEM analyses revealed an equiaxed recrystallized microstructure and dislocations rearranged to form high angle grain boundaries (HAGBs) upon dynamic recrystallization (DRX), respectively. The overall weak texture intensity was observed across the stir zone of FSPed samples due to the multiple passes. Incorporation of CNTs/SiC particles in Al5083 matrix resulted in the activation of Zener-Holloman mechanism and particlestimulated nucleation (PSN) mechanism by developing randomly oriented grains. In FSPed composites, SiC particles are dispersed homogeneously with good interfacial bonding and CNTs are partially reacted with an Al5083 matrix to form in-situ Al4C3 intermetallic compound. The maximum tensile strength of 361 MPa was obtained for Al5083-CNTs/SiC hybrid composite. The fracture surface of the SiC reinforced composite revealed that the voids initiation at the matrix-particle interface regions. © 2019 Elsevier B.V. All rights reserved.
Keywords: Friction stir processing SiC CNTs Mechanical properties Particle-stimulated nucleation
1. Introduction Aluminum metal matrix composites (AMMCs) have emerged as high-performance materials for a variety of engineering applications due to their high strength to weight ratio, high stiffness and wear resistance [1e4]. However, AMMCs prepared by conventional route shows poorer metallurgical and mechanical properties. Therefore, Friction stir processing (FSP) is a novel solid-state process for fabrication of AMMCs, which causes localized microstructural modification and enhances specific properties by controlling surface layers of the metallic materials [5e7]. The applications of FSP includes grain refinement [8], process surface composites [9], impart superplasticity [10], and microstructural modification of cast alloys [11]. Several research works on the fabrication of various AMMCs were reported in the literature [12e20]. Mishra et al. [12] studied the effect of shoulder position, and traverse speed on SiC reinforced Al5083 matrix composite. Their results showed that at optimum conditions, SiC particles were homogenous distributed in the aluminum matrix and two-fold increase in hardness was
* Corresponding author. E-mail address: [email protected] (S. Muthukumaran). https://doi.org/10.1016/j.jallcom.2019.05.232 0925-8388/© 2019 Elsevier B.V. All rights reserved.
obtained for Al5083-SiC composite. Guo et al. [13] evaluate the effect of multiple passes on grain refinement and mechanical properties of AA6061-Al2O3 nanocomposite. The results revealed that homogenous microstructure with uniform particle distribution and enhanced mechanical properties were obtained after the fourth pass when compared to unreinforced AA6061 alloy. Yuvaraj et al. [14] incorporated micro and nano-sized B4C particles in Al5083 matrix and assessed the effect of the number of passes on microstructure, mechanical and tribological properties. Results showed that one pass FSP leads to cluster or agglomeration of particles and subsequent FSP passes eliminate clusters and redistribute the particles in an aluminum matrix which further improves the mechanical and tribological properties. Liu et al. [15] successfully fabricated Al-CNTs nanocomposite by powder metallurgy route and subsequent FSP passes were applied to homogenize the composite which resulted in a significant enhancement in the mechanical properties. Palanivel et al. [16] studied the influence of BN nanoparticles on the A16082-TiB2 hybrid composite. They reported that uniform distribution of particles without any particle segregation, good interfacial bonding and incorporated BN nanoparticles in an aluminum matrix enhances the wear resistance. Amra et al. [17] produced Al5083-SiC/CeO2 mono and hybrid
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composites. They investigated the effect of volume fraction of SiC and CeO2 particles on microstructure, mechanical and corrosion properties. The results showed that optimum mechanical and corrosion properties were obtained for Al5083-SiC (25%)-CeO2 (75%) hybrid composite. Hosseini et al. [18] fabricated Al5083CNTs/CeO2 mono and hybrid composites. They reported that the volume ratio of CNTs (75%) and CeO2 (25%) showed the highest hardness and tensile strength whereas, CeO2 reinforced composite alone shows the enhanced corrosion resistance. Al6061-Al2O3/CNTs mono and hybrid nanocomposite were fabricated by Du et al. [19] via FSP. The yield strength of hybrid composites increased by 70% in comparison with FSPed Al6061alloy. Eskandari et al. [20] fabricated Al8026-TiB2/Al2O3 hybrid composite layers by multi-pass FSP and investigated the effect of the reinforcement on microstructure, hardness, tensile and wear properties. Likewise, they also reported that hybrid composite results in better mechanical and wear properties. The hybrid aluminum composite shows significant enhancement in mechanical and wear properties compared to mono-particle reinforced aluminum composites and alloys. To the best of the author's knowledge, little information is available as regards the use of CNTs and or microsized SiC particles either separately or in aggregated form to produced mono and hybrid composites through FSP. Silicon carbide (SiC) particles extensively used as reinforcement in aluminum matrix composites because of its higher hardness and elastic modulus. CNTs are another potential reinforcement which has exceptionally high strength of 30 GPa and modulus of 1 TPa. With these unique properties, CNTs is an ideal reinforcement in aluminum composites to enhance the mechanical properties. This study aims to investigate the effect of CNTs/SiC particles either individually or in aggregate form on the evolution of grain structure, texture, and mechanical properties. 2. Materials and methods Rolled Al5083 alloy plates (150 80 6 mm3) were used as starting material in this work. Table 1 shows the chemical composition of the Al5083 alloy. SiC particles (particle size between 5 and 20 mm) and CNTs (purity: 98%, outer diameter: 05e20 nm, 1e10 mm in length) were used as reinforcements. Fig. 1 shows the SEM micrographs of SiC particles and CNTs. A groove of 1 mm wide and 2.5 mm deep was machined at the center of plates using wire electric discharge machining (W-EDM) is shown in Fig. 2a. Reinforcements were closely packed into the grooves with various compositions by volume fractions such as 100%SiC, 50%SiC-50% CNTs and 100%CNTs are shown in Fig. 2b. The groove surface was closed by using pinless tool to void the particles from dispersing during FSP is shown in Fig. 2c. All the samples were subjected to three passes for the homogeneous dispersion of reinforcements with the rotational speed of 1600 rpm, and transverse speed of 20 mm/min are shown in Fig. 2d. FSP was also carried out on base alloy without reinforcement with the same experimental conditions for comparison purposes. Details of FSP machine, process parameters, and tool used to produce composites were discussed elsewhere [9,21]. The microstructural studies of the FSPed samples were characterized by optical microscopy, scanning electron microscopy equipped with energy-dispersive X-ray spectroscopy, electron backscattered diffraction (EBSD) and transmission electron microscopy (TEM). For EBSD studies, samples were electro-
Table 1 Chemical compositions of the Al5083 alloy. Elements
Mg
Si
Mn
Cr
Fe
Cu
Zn
Ti
AA
Wt.%
4.7
0.2
0.9
0.09
0.7
0.28
0.05
0.05
BAL.
83
polished using perchlorate alcohol solution at 12 C for 20 s. The electro-polished samples were then observed under FEI Quanta FEG-SEM equipped with a TSL-OIM software operating at 30 kV. The microhardness profile was obtained across the cross-section of FSPed samples using an indenter with a 100 g load for 15 s. TEM samples were sliced from the FSPed region and polished to a thickness of 80 mm, and then ion milled until perforation. The ion milled samples were then carefully observed under JEOL JEM 2100 HRTEM, operating at 200 keV. Tensile tests were carried out on Tinius Olsen (H10KT) tensile testing machine at a strain rate of 8.33 104s1. Further, the fracture surfaces were analyzed under SEM to understand the nature of the failure. 3. Results and discussion 3.1. Particle dispersion During the processing of composites via FSP, the primary issue is to incorporate a large volume fraction of reinforcements with better distribution at optimum processing conditions. It has been reported that the uniform distribution of particles with one pass is a very challenging task due to the complex material flow and intermixing mechanism imposed by tool [21,22]. It is believed that particle distribution can be improved by employing multiple passes. Fig. 3a shows the uniform distribution of SiC particles in the stir zone of Al5083-SiC composite. The distribution of SiC particles can be attributed to the vigorous rotation of the tool. During FSP, stirring action of tool probe breaks off the SiC particles into smaller fragments. The size of these fragments varies from few microns to nano in size as shown in Fig. 3b. Whereas in the case of CNTs reinforced Al5083 composite exhibits a large amount of CNTs clusters/agglomeration in the stir zone on the advancing side during the first pass and are not embedded in the matrix as shown in Fig. 3c. Therefore, to re-distribute CNTs more homogeneously in the stir zone, the second and third pass was carried out in reversed direction rather than in same direction (i.e., advancing side of first and third pass become the retreating side in the second pass) [23]. The results revealed that multi-pass improves the CNTs distribution homogeneously throughout the stir zone by reducing the aspect ratio of CNTs to the greater extent and thereby obtaining good interfacial bonding between Al5083 matrix and CNTs (Fig. 3d). Furthermore, FSP partially initiates an interfacial reaction between Al5083 matrix and CNTs to form in-situ Al4C3 intermetallic compound as supported with TEM results (Fig. 8d and e). 3.2. Grain structure evolution Fig. 4 shows the EBSD micrograph (IPF þ grain boundary) and grain size distribution of the Al5083 alloy. The base Al5083 alloy comprises of elongated grains parallel to the rolling direction as shown in Fig. 4a. The average grain size of Al5083 alloy is about 49 mm as shown in Fig. 4b. Table 2 shows the summarised EBSD results of Al5083 alloy and FSPed samples. The fraction of high angle grain boundaries (HAGBs) and low angle grain boundaries (LAGBs) in Al5083 alloy are 10% and 90%, respectively. After three passes, FSPed samples with and without reinforcement shows fine equiaxed recrystallized grains structure (Fig. 5aed). From Table 2, it can also be seen that FSPed samples have a higher fraction of HAGBs compared to base Al5083 alloy. Furthermore, Fig. 6 shows the grain size distribution of FSPed samples. The average grain size of FSPed samples of Al5083 alloy, Al5083-SiC, Al5083-SiC/CNTs, and Al5083-CNTs is 5.83 mm, 5.24 mm, 6.06 mm, and 6.36 mm respectively. Fig. 7 shows the Al5083-SiC/CNTs image quality map superimposed with the grain boundary map obtained from EBSD. Two types of low angle boundaries were highlighted in Fig. 7a and b
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Fig. 1. SEM micrographs of (a) SiC particles and (b) CNTs.
Fig. 2. Schematic view of different stages of FSP (a) Groove prepared on the plate surface, (b) Reinforcements are incorporated in the groove, (c) Pinless tool was employed to close the top surface of the groove, and (d) Multiple FSP passes were applied for the homogeneous distribution of reinforcements.
(red, 2 e5 and green, 5 e15 ). Second, boundaries beyond (15 ) are called a high angle boundary, which is indicated in blue color (Fig. 7a and b). It is considered that various mechanism are associated with grain structure evolution during FSP which includes; dynamic recovery (DRV), geometric dynamic recrystallization (GDRX), continuous dynamic recrystallization (CDRX) and discontinuous dynamic recrystallization (DDRX). During FSP, materials are severely deformed at high temperature (i.e., below its melting point) and high strain. DRV quickly occurs in aluminum owing to its
high stacking fault energy (SFE) which results in increases in flow stress, dislocation grows and annihilate [24e27]. However, the recovery rate increases with dislocation density. These dislocations begin to rearrange to form subgrains boundaries (2 e5 ). Upon dynamic equilibrium, subgrain coalescence begins to form by absorbing more and more dislocations and transforms into low angle grain boundaries (5 e15 ) are shown in white arrows in Fig. 7b. Therefore, when subgrains reach a critical size, CDRX begins and slowly rotate the subgrains to increase their misorientation and
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Fig. 3. (a) Optical micrograph of Al5083-SiC composite showing even distribution of SiC particles after three passes FSP in stir zone area (b) SEM micrograph showing fragmented SiC particles after three pass FSP, (c) CNTs agglomeration after the first pass, (d) Even dispersion of CNTs after third pass FSP.
Fig. 4. (a) EBSD map (IPF þ grain boundary) and, (b) grain size distribution of the Al5083 alloy.
Table 2 Summarised results of grain size, LAGBs, HAGBs for various materials. Samples
Al5083 FSPed Al5083 Al5083-SiC Al5083-SiC/CNT's Al5083-CNT's
Grain size (mm)
40.00 5.83 5.24 6.06 6.36
Fraction of low angle grain boundaries (fLAGBs) in %
Fraction of high angle grain boundaries (fHAGBs) in %
(2e5 )
(5e15 )
(15e180 )
69.9 12.8 16.4 19.0 16.8
19.7 13.5 15.1 17.1 15.2
10.4 73.7 68.5 63.8 67.9
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Fig. 5. EBSD map (IPF þ grain boundary) of FSPed samples (a) Al5083 alloy, (b) Al5083-CNTs, (c) Al5083-SiC/CNTs, (d) Al5083-SiC.
Fig. 6. Grain size distribution of FSPed samples.
transform them from LAGBs to HAGBS (>15 ) are illustrated in white arrows in Fig. 7b. It has been found that subgrain size often depends upon the Zener-Hollomon parameter, i.e., Z ¼ ℇ exp (Q/RT) where ℇ - strain rate, Q - activation energy, R - gas constant, and T temperature. Other than DRV and CDRX, two additional
mechanisms involved in grain refinement of FSPed samples are particle-stimulated nucleation (PSN) and Zener pinning. During FSP/FSW, second phase particles are refined and redistributed throughout the stir zone because the solvus temperature of these particles is higher than the solidus temperature of Al5083 alloy. Many small particles of Al6(Mn, Fe) were observed in the stir zone of the friction stir welded Al1080 and Al5083 alloy [28]. In the present case, the occurrence of DDRX mechanism does not observe due to the presence of second phase particles which can effectively pin the grains and subgrain boundaries to prevent their migration and grain growth. Khodabakhshi et al. [29] reported two types of large complex particles such as (Fe, Mn, Cr)3SiAl12 and small Mg2Si in Al5052 alloy. Their results indicated that after the thermomechanical process, the morphology of particles was changed to elongated cuboidal (Al-Fe-Mn-Si) and spherical (Mg2Si) particles. Moreover, the sheared SiC particles also hinder the grain growth leading to fine equiaxed microstructure. The presence of fine precipitates and fragmented SiC particles in the Al5083 matrix pronounces the occurrence of PSN assisted CDRX as a dominant mechanism for grain refinement in the stir zone. Fig. 8 shows the bright field TEM micrographs of Al5083-SiC/ CNTs. The fine second phase particles are visible in the stir zone (Fig. 8a). It was observed that due to repeated passes, large sized second phase particles are refined and redistributed homogeneously throughout the stir zone. Fig. 8b shows the arrangement of dislocations into subgrain boundaries (white arrows). These dislocations are merged and gradually increases its misorientation
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Fig. 7. (a) Grain boundary map superimposed on image quality of the Al5083-SiC/CNTs, boundaries, (b) boundaries with mixed orientation showing LAGBS and HAGBs.
leading to the formation of subgrain boundary. Fig. 8c shows the second phase particles hinders the dislocation motions. These particles are incoherent and can pin the grain boundaries. However, these particles also act as nucleation sites for recrystallization. These hindrances generate geometrically necessary dislocations, which further strengthen the matrix by PSN mechanism. Fig. 8d shows the survived CNTs during FSP. Formation of in-situ Al4C3 intermetallic compound on interface and middle part of twisted CNTs is mainly due to the repeated FSP passes and high strain. However, the full reaction was not observed between the CNTs and the Al matrix. The interfacial reaction always occurs on the surface of CNTs due to its amorphous carbon film in the order of ~1 nm which offer preferred sites for reaction. Interfacial Al4C3 intermetallic compound was well bonded with CNTs and Al matrix, seen in Fig. 8e.
3.3. Crystallographic texture It is well known that FSP of face-centered cubic (FCC) metals exhibit a predominant simple shear deformation texture [30,31]. Fig. 9a shows the ideal simple shear deformation texture of FCC metals on a {111} pole figure and position of ideal orientations are also labeled. For comparative purposes, {111} pole figure of Al5083CNTs, Al5083-SiC, Al5083-SiC-CNTs, and FSPed Al5083 alloy are shown in Fig. 9cef respectively. The discrepancies between the ideal and experimental texture components are mainly due to 3 tool tilt angle and threads on tool probe which may cause a slight variation in shear plane and direction. The EBSD map (Fig. 4a) and pole figure (Fig. 9b) of base Al5083 alloy clearly shows an elongated grain structure and strong rolling texture dominated by the S components and with the minor presence of the B component with a peak intensity of 5.75. As viewed in Fig. 9cef, strong {112} <110> B/B texture component for FSPed Al5083-SiC, {111} <112> A*1/A*2 texture and {112} <110> B/B texture component for FSPed Al5083CNTs, and {001} <110> C texture, {111} <112> A*1/A*2 texture for FSPed Al5083-SiC/CNTs and {001} <110> C texture, {111} <112>A*2 texture and {111} <110> A texture components for FSPed Al5083 alloy were observed in the stir zone respectively. Addition of the reinforcements and re-precipitation of second phase particles weakened the texture due to activation of PSN mechanism and Zener pinning. Furthermore, other ideal shear component are also present to some extent in all FSPed samples. The variation in texture components in the FSPed samples is mainly due to the presence of reinforcements and their distributional effect, because the process parameters is the same for all FSPed samples. It is reported that at the lower strain value, A/A texture component is
dominant and rapidly switched to C texture component as the strain value reached between 3 and 5 and beyond which B/B texture component slowly replaces the C texture component [32,33]. 3.4. Microhardness Microhardness profiles of the FSPed regions of composites along the cross-section after three FSP passes are shown in Fig. 10. The average hardness of the base Al5083 alloy was found to be 79.25 ± 3.96 HV. It can be observed that the average hardness of FSPed Al5083 alloy was about 94.00 ± 4.70 HV, which is higher when compared with base Al5083 alloy. The factors which are responsible for the improved hardness of FSPed Al5083 alloy are grain refinement and re-precipitation of second phase particles. Initially, base Al5083 alloy has elongated grains with an average size of 49 mm (Fig. 4a) and large complex shaped second phase particles which are then replaced by fine equiaxed grain structure, fragmentation, and re-distribution of precipitates throughout the stir zone by the thermo-mechanical process of FSP. However, these second phase particles do not dissolve in the stir zone. Thus only fragmentation of particles occurs [28,29]. The aluminum 5083 alloys reinforced with SiC and CNTs either individually or in aggregate form shows an improvement in the hardness compared to FSPed Al5083 alloy and base Al5083 alloy. The average hardness of Al5083-SiC, Al5083-SiC/CNTs and Al5083-CNTs are 124.50 ± 6.22 HV, 112.50 ± 5.62 HV, and 107.50 ± 5.37 HV respectively, which is 1.5, 1.4 and 1.35 times, higher than base Al5083 alloy. The enhanced hardness is mainly due to a homogeneous distribution of CNTs and SiC particle results in dislocations pinning and impeding grain growth, good interfacial bonding between the matrix and particles. 3.5. Tensile properties Fig. 11 depicts the engineering stress-strain curves for Al5083 alloy and FSPed samples. Tensile tests were carried out to determine the tensile strength, yield strength, and elongation of FSPed samples and base Al5083 alloy. It is evident that all the FSPed composite samples exhibited higher tensile strength with a subsequent drop in the elongation as compared to base Al5083 alloy. Among the FSPed samples, SiC/CNTs reinforced Al5083 matrix composites shows the maximum tensile strength of 361 MPa. Whereas Al5083-SiC, Al5083-CNTs, FSPed Al5083, and Al5083 alloy shows the ultimate tensile strength of 335 MPa, 316 MPa, 306 MPa, and 298 MPa respectively. However, FSPed Al5083 alloy shows higher elongation compared to other FSPed samples and base
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Fig. 8. Bright field TEM images of Al5083-SiC/CNTs showing (a) Fine second phase particles in stir zone (b) dislocation rearrangement into subgrain boundaries (white arrows) (c) second phase fine particle obstruct dislocation motion, (d) survived twisted CNTs with Al4C3 phase (e) interface between CNTs and Al matrix leading to Al4C3 phase formation.
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Fig. 9. {111} pole figure of (a) ideal orientations of FCC metals under simple shear deformation texture with fiber component labeled [30,31], (b) Al5083 alloy, (c) FSPed Al5083-SiC, (d) FSPed Al5083-CNTs, (e) FSPed Al5083-SiC/CNTs and (f) FSPed Al5083 alloy.
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Hosseini et al. [18] is having good agreement with our results. Similar results of Al5083-SiC composites were reported by Amra et al. [17]. However, in this work, Al5083-CNTs/SiC composites show a higher ultimate tensile strength of 361 MPa which clearly shows that FSP is effectively utilized to fabricate AMMCs without any detrimental products. The properties difference mainly depends upon the chemical composition, processing parameters, and morphology of reinforcement used. 3.6. Fractography
Fig. 10. Microhardness profile across the transverse section of the Al5083 alloy and FSPed samples.
Fig. 12 depicts the SEM fracture surface of Al5083 alloy and FSPed samples. Fig. 12a shows the presence of larger dimples and voids indicating a typical ductile mode of failure. In the FSPed Al5083 alloy (Fig. 12b), a similar type of ductile failure with marginally larger dimples and voids are observed which indicates that FSPed Al5083 alloy exhibit higher elongation before fracture as compared to base Al5083 alloy and other FSPed composite samples. The improved strength and elongation of FSPed Al5083 alloy are attributed to grain refinement owing to dynamic recrystallization [17]. Some of the second phase intermetallic particles are present in the core of the dimples of Al5083 alloy and FSPed Al5083 alloy. The FSPed composite samples (Fig. 12c and d) show the bimodal distribution of smaller to larger dimples and voids are attributed to grain refinement and homogenous distribution of reinforcement [18,34]. During tensile loading, voids are formed around the particles and mobility of dislocation hinder by reinforcement which further increases stress to deform the material. Failure occurs due coalescence of voids, followed by growth. The fracture sites in Al5083-SiC composite sample shows that the SiC particles found at the bottom of the dimples are shown in Fig. 12d. The fracture occurs at particle/matrix interfaces showing good cohesion between the particle and matrix resulted in sound interfacial bonding. 4. Conclusions In the present study, the effect of CNTs/SiC either separately or in aggregate form, on the microstructure, texture evolution, and mechanical properties of Al5083 composites were investigated. Based on the experimental results, the following conclusions were drawn:
Fig. 11. Stress-strain curves of FSPed samples and base Al5083 alloy.
Al5083 alloy. Al5083 matrix strengthening was obtained by grain refinement, dispersion of particles and re-precipitates, and dislocations. Uniform dispersion of reinforcement, clean interface, and strong interfacial bonding enhances the dispersion strengthening. Second, fine-grained material offers more resistance to deformation and substantial strengthening by grain refining. Third, due to the difference in coefficient of thermal expansion between reinforcement and matrix induces additional dislocations during material flow. The mobility of dislocations is hinder by the reinforcement, which further enhances the strength of materials. The tensile properties of the present work are comparable with values reported in the literature. Bauri et al. [34] have reported tensile strength of 296 MPa and 337 MPa, and % elongation of 25% and 33% for base A15083 alloy and FSPed Al5083 alloy respectively. The work of Yuvaraj et al. [35] having a close agreement with present work, i.e., tensile strength and % elongation of Al5083 alloy before and after FSP are 300 MPa, 27.1%, and 312 MPa, 30.2% respectively. Also, CNTs reinforced with Al5083 alloy reported by
1. CNTs/SiC reinforced Al5083 composites are successfully processed using multi-pass FSP with a uniform distribution of SiC and CNTs. FSP also results in significant grain refinement, fragmentation, and re-precipitation of second phase particles and partially initiated an interfacial reaction between the Al matrix and CNTs to form in-situ Al4C3 intermetallic compound. 2. The dominant texture component observed in the stir zone of FSPed samples are {112} <110> B/B for Al5083-SiC, {111} <112> A*1/A*2 texture and {112} <110> B/B texture for Al5083-CNTs, {001} <110> C texture and {111} <112> A*1/A*2 texture for Al5083-SiC/CNTs. 3. Incorporation of CNTs/SiC and second phase particles resulted in the activation of Zener-Holloman mechanism and particlestimulated nucleation (PSN) mechanism by developing randomly oriented grains. 4. The average hardness of Al5083-SiC, Al5083-SiC/CNTs and Al5083-CNTs are 124.50 ± 6.22 HV, 112.50 ± 5.62 HV, and 107.50 ± 5.37 HV respectively, which is 1.5, 1.4 and 1.35 times, higher hardness than base Al5083 alloy. 5. The maximum tensile strength of 361 MPa was obtained for Al5083-SiC/CNTs hybrid composite whereas, Al5083-SiC, Al5083-CNTs, FSPed Al5083 alloy and base Al5083 alloy shows
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Fig. 12. Fracture surface (a) Al5083, (b) FSPed Al5083 alloy, (c) Al5083-SiC/CNTs (d) Al5083-CNTs, (e) Al5083-SiC.
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the tensile strength of 335 MPa, 316 MPa, 306 MPa, and 298 MPa, respectively. [17]
Acknowledgments The author (Vikram Kumar S. Jain), acknowledge Department of Science & Technology, Govt. of India for sponsoring him to pursue Ph.D. under INSPIRE Fellowship (DST/INSPIRE Fellowship/2015/ IF150488). The authors acknowledge the use of the National Facility of Texture & OIM, IIT Bombay for their support in carrying out the EBSD studies. The authors thank Dr. Devinder Yadav, Assistant Professor, IIT Patna for a useful discussion of the manuscript.
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nanoparticles on microstructure and wear behavior of AA6082/TiB2 hybrid aluminum composites synthesized by friction stir processing, Mater. Des. 106 (2016) 195e204. M. Amra, K. Ranjbar, R. Dehmolaei, Mechanical properties and corrosion behavior of CeO2 and SiC incorporated Al5083 alloy surface composites, J. Mater. Eng. Perform. 24 (2015) 3169e3179. S.A. Hosseini, K. Ranjbar, R. Dehmolaei, A.R. Amirani, Fabrication of Al5083 surface composites reinforced by CNTs and cerium oxide nano particles via friction stir processing, J. Alloys Compd. 622 (2015) 725e733. Z. Du, M. Jen, J. Feng, G. Bi, J. Wei, Fabrication of a new Al-Al2O3 -CNTs composite using friction stir processing (FSP), Mater. Sci. Eng. A 667 (2016) 125e131. H. Eskandari, R. Taheri, F. Khodabakhshi, Friction-stir processing of an AA8026-TiB2-Al2O3 hybrid nanocomposite: microstructural developments and mechanical properties, Mater. Sci. Eng. A 660 (2016) 84e96. V.K.S. Jain, V. James, S. Muthukumaran, Effect of first and second passes on microstructure and wear properties of titanium dioxide-reinforced aluminum surface composite via friction stir processing, Arabian J. Sci. Eng. 44 (2019) 949e957. R. Bauri, D. Yadav, G. Suhas, Effect of friction stir processing (FSP) on microstructure and properties of Al-TiC in situ composite, Mater. Sci. Eng. A 528 (2011) 4732e4739. E.R.I. Mahmoud, M. Takahashi, T. Shibayanagi, K. Ikeuchi, Effect of friction stir processing tool probe on fabrication of SiC particle reinforced composite on aluminium surface, Sci. Technol. Weld. Join. 14 (2009) 413e425. A. Rollett, F. Humphreys, G.S. Rohrer, M. Hatherly, Recrystallization and Related Annealing Phenomena, second ed., Pergamon, 2004, pp. 415e628. T.R. McNelley, S. Swaminathan, J.Q. Su, Recrystallization mechanisms during friction stir welding/processing of aluminum alloys, Scripta Mater. 58 (2008) 349e354. J.Q. Su, T.W. Nelson, T.R. McNelley, R.S. Mishra, Development of nanocrystalline structure in Cu during friction stir processing (FSP), Mater. Sci. Eng. A 528 (2011) 5458e5464. J.Q. Su, T.W. Nelson, C.J. Sterling, Microstructure evolution during FSW/FSP of high strength aluminum alloys, Mater. Sci. Eng. A 405 (2005) 277e286. Y.S. Sato, S.H.C. Park, H. Kokawa, Microstructural factors governing hardness in friction-stir welds of solid-solution-hardened Al alloys, Metall. Mater. Trans. A 32 (2001) 3033e3042. F. Khodabakhshi, A.P. Gerlich, P. Svec, Reactive friction-stir processing of an Al-Mg alloy with introducing multi-walled carbon nano-tubes (MW-CNTs): microstructural characteristics and mechanical properties, Mater. Char. 131 (2017) 359e373. Saiyi Li, Irene J. Beyerlein, Mark A.M. Bourke, Texture formation during equal channel angular extrusion of fcc and bcc materials: comparison with simple shear, Mater. Sci. Eng. A 394 (2005) 66e77. U.F.H.R. Suhuddin, S. Mironov, Y.S. Sato, H. Kokawa, Grain structure and texture evolution during friction stir welding of thin 6016 aluminum alloy sheets, Mater. Sci. Eng. A 527 (2010) 1962e1969. R.W. Fonda, J.F. Bingert, Texture variations in an aluminum friction stir weld, Scripta Mater. 57 (2007) 1052e1055. M.M.Z. Ahmed, B.P. Wynne, M.M. El-Sayed Seleman, W.M. Rainforth, A comparison of crystallographic texture and grain structure development in aluminum generated by friction stir welding and high strain torsion, Mater. Des. 103 (2016) 259e267. R. Bauri, D. Yadav, C.N. Shyam Kumar, B. Balaji, Tungsten particle reinforced Al 5083 composite with high strength and ductility, Mater. Sci. Eng. A 620 (2014) 67e75. N. Yuvaraj, S. Aravindan, Vipin, Wear characteristics of Al5083 surface hybrid nano-composites by friction stir processing, Trans. Indian Inst. Met. 70 (2016) 1111e1129.
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Development and characterization of Al5083-CNTs/SiC composites via friction stir processing Vikram Kumar S. Jain a, K.U. Yazar b, S. Muthukumaran a, * a b
Department of Metallurgical and Materials Engineering, National Institute of Technology, Tiruchirappalli, Tamil Nadu, 620 015, India Department of Materials Engineering, Indian Institute of Science, Bengaluru, 560012, India
a r t i c l e i n f o
a b s t r a c t
Article history: Received 7 March 2019 Received in revised form 26 April 2019 Accepted 19 May 2019 Available online 21 May 2019
Carbon nanotubes (CNTs) and or micron-sized silicon carbide (SiC) particles were reinforced with Al5083 matrix to develop mono and hybrid composites via friction stir processing (FSP). The effect of CNTs/SiC either individually or in aggregate form, on microstructural evolution, texture, and mechanical properties of friction stir processed (FSPed) Al5083 composites were studied. EBSD and TEM analyses revealed an equiaxed recrystallized microstructure and dislocations rearranged to form high angle grain boundaries (HAGBs) upon dynamic recrystallization (DRX), respectively. The overall weak texture intensity was observed across the stir zone of FSPed samples due to the multiple passes. Incorporation of CNTs/SiC particles in Al5083 matrix resulted in the activation of Zener-Holloman mechanism and particlestimulated nucleation (PSN) mechanism by developing randomly oriented grains. In FSPed composites, SiC particles are dispersed homogeneously with good interfacial bonding and CNTs are partially reacted with an Al5083 matrix to form in-situ Al4C3 intermetallic compound. The maximum tensile strength of 361 MPa was obtained for Al5083-CNTs/SiC hybrid composite. The fracture surface of the SiC reinforced composite revealed that the voids initiation at the matrix-particle interface regions. © 2019 Elsevier B.V. All rights reserved.
Keywords: Friction stir processing SiC CNTs Mechanical properties Particle-stimulated nucleation
1. Introduction Aluminum metal matrix composites (AMMCs) have emerged as high-performance materials for a variety of engineering applications due to their high strength to weight ratio, high stiffness and wear resistance [1e4]. However, AMMCs prepared by conventional route shows poorer metallurgical and mechanical properties. Therefore, Friction stir processing (FSP) is a novel solid-state process for fabrication of AMMCs, which causes localized microstructural modification and enhances specific properties by controlling surface layers of the metallic materials [5e7]. The applications of FSP includes grain refinement [8], process surface composites [9], impart superplasticity [10], and microstructural modification of cast alloys [11]. Several research works on the fabrication of various AMMCs were reported in the literature [12e20]. Mishra et al. [12] studied the effect of shoulder position, and traverse speed on SiC reinforced Al5083 matrix composite. Their results showed that at optimum conditions, SiC particles were homogenous distributed in the aluminum matrix and two-fold increase in hardness was
* Corresponding author. E-mail address: [email protected] (S. Muthukumaran). https://doi.org/10.1016/j.jallcom.2019.05.232 0925-8388/© 2019 Elsevier B.V. All rights reserved.
obtained for Al5083-SiC composite. Guo et al. [13] evaluate the effect of multiple passes on grain refinement and mechanical properties of AA6061-Al2O3 nanocomposite. The results revealed that homogenous microstructure with uniform particle distribution and enhanced mechanical properties were obtained after the fourth pass when compared to unreinforced AA6061 alloy. Yuvaraj et al. [14] incorporated micro and nano-sized B4C particles in Al5083 matrix and assessed the effect of the number of passes on microstructure, mechanical and tribological properties. Results showed that one pass FSP leads to cluster or agglomeration of particles and subsequent FSP passes eliminate clusters and redistribute the particles in an aluminum matrix which further improves the mechanical and tribological properties. Liu et al. [15] successfully fabricated Al-CNTs nanocomposite by powder metallurgy route and subsequent FSP passes were applied to homogenize the composite which resulted in a significant enhancement in the mechanical properties. Palanivel et al. [16] studied the influence of BN nanoparticles on the A16082-TiB2 hybrid composite. They reported that uniform distribution of particles without any particle segregation, good interfacial bonding and incorporated BN nanoparticles in an aluminum matrix enhances the wear resistance. Amra et al. [17] produced Al5083-SiC/CeO2 mono and hybrid
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composites. They investigated the effect of volume fraction of SiC and CeO2 particles on microstructure, mechanical and corrosion properties. The results showed that optimum mechanical and corrosion properties were obtained for Al5083-SiC (25%)-CeO2 (75%) hybrid composite. Hosseini et al. [18] fabricated Al5083CNTs/CeO2 mono and hybrid composites. They reported that the volume ratio of CNTs (75%) and CeO2 (25%) showed the highest hardness and tensile strength whereas, CeO2 reinforced composite alone shows the enhanced corrosion resistance. Al6061-Al2O3/CNTs mono and hybrid nanocomposite were fabricated by Du et al. [19] via FSP. The yield strength of hybrid composites increased by 70% in comparison with FSPed Al6061alloy. Eskandari et al. [20] fabricated Al8026-TiB2/Al2O3 hybrid composite layers by multi-pass FSP and investigated the effect of the reinforcement on microstructure, hardness, tensile and wear properties. Likewise, they also reported that hybrid composite results in better mechanical and wear properties. The hybrid aluminum composite shows significant enhancement in mechanical and wear properties compared to mono-particle reinforced aluminum composites and alloys. To the best of the author's knowledge, little information is available as regards the use of CNTs and or microsized SiC particles either separately or in aggregated form to produced mono and hybrid composites through FSP. Silicon carbide (SiC) particles extensively used as reinforcement in aluminum matrix composites because of its higher hardness and elastic modulus. CNTs are another potential reinforcement which has exceptionally high strength of 30 GPa and modulus of 1 TPa. With these unique properties, CNTs is an ideal reinforcement in aluminum composites to enhance the mechanical properties. This study aims to investigate the effect of CNTs/SiC particles either individually or in aggregate form on the evolution of grain structure, texture, and mechanical properties. 2. Materials and methods Rolled Al5083 alloy plates (150 80 6 mm3) were used as starting material in this work. Table 1 shows the chemical composition of the Al5083 alloy. SiC particles (particle size between 5 and 20 mm) and CNTs (purity: 98%, outer diameter: 05e20 nm, 1e10 mm in length) were used as reinforcements. Fig. 1 shows the SEM micrographs of SiC particles and CNTs. A groove of 1 mm wide and 2.5 mm deep was machined at the center of plates using wire electric discharge machining (W-EDM) is shown in Fig. 2a. Reinforcements were closely packed into the grooves with various compositions by volume fractions such as 100%SiC, 50%SiC-50% CNTs and 100%CNTs are shown in Fig. 2b. The groove surface was closed by using pinless tool to void the particles from dispersing during FSP is shown in Fig. 2c. All the samples were subjected to three passes for the homogeneous dispersion of reinforcements with the rotational speed of 1600 rpm, and transverse speed of 20 mm/min are shown in Fig. 2d. FSP was also carried out on base alloy without reinforcement with the same experimental conditions for comparison purposes. Details of FSP machine, process parameters, and tool used to produce composites were discussed elsewhere [9,21]. The microstructural studies of the FSPed samples were characterized by optical microscopy, scanning electron microscopy equipped with energy-dispersive X-ray spectroscopy, electron backscattered diffraction (EBSD) and transmission electron microscopy (TEM). For EBSD studies, samples were electro-
Table 1 Chemical compositions of the Al5083 alloy. Elements
Mg
Si
Mn
Cr
Fe
Cu
Zn
Ti
AA
Wt.%
4.7
0.2
0.9
0.09
0.7
0.28
0.05
0.05
BAL.
83
polished using perchlorate alcohol solution at 12 C for 20 s. The electro-polished samples were then observed under FEI Quanta FEG-SEM equipped with a TSL-OIM software operating at 30 kV. The microhardness profile was obtained across the cross-section of FSPed samples using an indenter with a 100 g load for 15 s. TEM samples were sliced from the FSPed region and polished to a thickness of 80 mm, and then ion milled until perforation. The ion milled samples were then carefully observed under JEOL JEM 2100 HRTEM, operating at 200 keV. Tensile tests were carried out on Tinius Olsen (H10KT) tensile testing machine at a strain rate of 8.33 104s1. Further, the fracture surfaces were analyzed under SEM to understand the nature of the failure. 3. Results and discussion 3.1. Particle dispersion During the processing of composites via FSP, the primary issue is to incorporate a large volume fraction of reinforcements with better distribution at optimum processing conditions. It has been reported that the uniform distribution of particles with one pass is a very challenging task due to the complex material flow and intermixing mechanism imposed by tool [21,22]. It is believed that particle distribution can be improved by employing multiple passes. Fig. 3a shows the uniform distribution of SiC particles in the stir zone of Al5083-SiC composite. The distribution of SiC particles can be attributed to the vigorous rotation of the tool. During FSP, stirring action of tool probe breaks off the SiC particles into smaller fragments. The size of these fragments varies from few microns to nano in size as shown in Fig. 3b. Whereas in the case of CNTs reinforced Al5083 composite exhibits a large amount of CNTs clusters/agglomeration in the stir zone on the advancing side during the first pass and are not embedded in the matrix as shown in Fig. 3c. Therefore, to re-distribute CNTs more homogeneously in the stir zone, the second and third pass was carried out in reversed direction rather than in same direction (i.e., advancing side of first and third pass become the retreating side in the second pass) [23]. The results revealed that multi-pass improves the CNTs distribution homogeneously throughout the stir zone by reducing the aspect ratio of CNTs to the greater extent and thereby obtaining good interfacial bonding between Al5083 matrix and CNTs (Fig. 3d). Furthermore, FSP partially initiates an interfacial reaction between Al5083 matrix and CNTs to form in-situ Al4C3 intermetallic compound as supported with TEM results (Fig. 8d and e). 3.2. Grain structure evolution Fig. 4 shows the EBSD micrograph (IPF þ grain boundary) and grain size distribution of the Al5083 alloy. The base Al5083 alloy comprises of elongated grains parallel to the rolling direction as shown in Fig. 4a. The average grain size of Al5083 alloy is about 49 mm as shown in Fig. 4b. Table 2 shows the summarised EBSD results of Al5083 alloy and FSPed samples. The fraction of high angle grain boundaries (HAGBs) and low angle grain boundaries (LAGBs) in Al5083 alloy are 10% and 90%, respectively. After three passes, FSPed samples with and without reinforcement shows fine equiaxed recrystallized grains structure (Fig. 5aed). From Table 2, it can also be seen that FSPed samples have a higher fraction of HAGBs compared to base Al5083 alloy. Furthermore, Fig. 6 shows the grain size distribution of FSPed samples. The average grain size of FSPed samples of Al5083 alloy, Al5083-SiC, Al5083-SiC/CNTs, and Al5083-CNTs is 5.83 mm, 5.24 mm, 6.06 mm, and 6.36 mm respectively. Fig. 7 shows the Al5083-SiC/CNTs image quality map superimposed with the grain boundary map obtained from EBSD. Two types of low angle boundaries were highlighted in Fig. 7a and b
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Fig. 1. SEM micrographs of (a) SiC particles and (b) CNTs.
Fig. 2. Schematic view of different stages of FSP (a) Groove prepared on the plate surface, (b) Reinforcements are incorporated in the groove, (c) Pinless tool was employed to close the top surface of the groove, and (d) Multiple FSP passes were applied for the homogeneous distribution of reinforcements.
(red, 2 e5 and green, 5 e15 ). Second, boundaries beyond (15 ) are called a high angle boundary, which is indicated in blue color (Fig. 7a and b). It is considered that various mechanism are associated with grain structure evolution during FSP which includes; dynamic recovery (DRV), geometric dynamic recrystallization (GDRX), continuous dynamic recrystallization (CDRX) and discontinuous dynamic recrystallization (DDRX). During FSP, materials are severely deformed at high temperature (i.e., below its melting point) and high strain. DRV quickly occurs in aluminum owing to its
high stacking fault energy (SFE) which results in increases in flow stress, dislocation grows and annihilate [24e27]. However, the recovery rate increases with dislocation density. These dislocations begin to rearrange to form subgrains boundaries (2 e5 ). Upon dynamic equilibrium, subgrain coalescence begins to form by absorbing more and more dislocations and transforms into low angle grain boundaries (5 e15 ) are shown in white arrows in Fig. 7b. Therefore, when subgrains reach a critical size, CDRX begins and slowly rotate the subgrains to increase their misorientation and
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Fig. 3. (a) Optical micrograph of Al5083-SiC composite showing even distribution of SiC particles after three passes FSP in stir zone area (b) SEM micrograph showing fragmented SiC particles after three pass FSP, (c) CNTs agglomeration after the first pass, (d) Even dispersion of CNTs after third pass FSP.
Fig. 4. (a) EBSD map (IPF þ grain boundary) and, (b) grain size distribution of the Al5083 alloy.
Table 2 Summarised results of grain size, LAGBs, HAGBs for various materials. Samples
Al5083 FSPed Al5083 Al5083-SiC Al5083-SiC/CNT's Al5083-CNT's
Grain size (mm)
40.00 5.83 5.24 6.06 6.36
Fraction of low angle grain boundaries (fLAGBs) in %
Fraction of high angle grain boundaries (fHAGBs) in %
(2e5 )
(5e15 )
(15e180 )
69.9 12.8 16.4 19.0 16.8
19.7 13.5 15.1 17.1 15.2
10.4 73.7 68.5 63.8 67.9
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Fig. 5. EBSD map (IPF þ grain boundary) of FSPed samples (a) Al5083 alloy, (b) Al5083-CNTs, (c) Al5083-SiC/CNTs, (d) Al5083-SiC.
Fig. 6. Grain size distribution of FSPed samples.
transform them from LAGBs to HAGBS (>15 ) are illustrated in white arrows in Fig. 7b. It has been found that subgrain size often depends upon the Zener-Hollomon parameter, i.e., Z ¼ ℇ exp (Q/RT) where ℇ - strain rate, Q - activation energy, R - gas constant, and T temperature. Other than DRV and CDRX, two additional
mechanisms involved in grain refinement of FSPed samples are particle-stimulated nucleation (PSN) and Zener pinning. During FSP/FSW, second phase particles are refined and redistributed throughout the stir zone because the solvus temperature of these particles is higher than the solidus temperature of Al5083 alloy. Many small particles of Al6(Mn, Fe) were observed in the stir zone of the friction stir welded Al1080 and Al5083 alloy [28]. In the present case, the occurrence of DDRX mechanism does not observe due to the presence of second phase particles which can effectively pin the grains and subgrain boundaries to prevent their migration and grain growth. Khodabakhshi et al. [29] reported two types of large complex particles such as (Fe, Mn, Cr)3SiAl12 and small Mg2Si in Al5052 alloy. Their results indicated that after the thermomechanical process, the morphology of particles was changed to elongated cuboidal (Al-Fe-Mn-Si) and spherical (Mg2Si) particles. Moreover, the sheared SiC particles also hinder the grain growth leading to fine equiaxed microstructure. The presence of fine precipitates and fragmented SiC particles in the Al5083 matrix pronounces the occurrence of PSN assisted CDRX as a dominant mechanism for grain refinement in the stir zone. Fig. 8 shows the bright field TEM micrographs of Al5083-SiC/ CNTs. The fine second phase particles are visible in the stir zone (Fig. 8a). It was observed that due to repeated passes, large sized second phase particles are refined and redistributed homogeneously throughout the stir zone. Fig. 8b shows the arrangement of dislocations into subgrain boundaries (white arrows). These dislocations are merged and gradually increases its misorientation
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Fig. 7. (a) Grain boundary map superimposed on image quality of the Al5083-SiC/CNTs, boundaries, (b) boundaries with mixed orientation showing LAGBS and HAGBs.
leading to the formation of subgrain boundary. Fig. 8c shows the second phase particles hinders the dislocation motions. These particles are incoherent and can pin the grain boundaries. However, these particles also act as nucleation sites for recrystallization. These hindrances generate geometrically necessary dislocations, which further strengthen the matrix by PSN mechanism. Fig. 8d shows the survived CNTs during FSP. Formation of in-situ Al4C3 intermetallic compound on interface and middle part of twisted CNTs is mainly due to the repeated FSP passes and high strain. However, the full reaction was not observed between the CNTs and the Al matrix. The interfacial reaction always occurs on the surface of CNTs due to its amorphous carbon film in the order of ~1 nm which offer preferred sites for reaction. Interfacial Al4C3 intermetallic compound was well bonded with CNTs and Al matrix, seen in Fig. 8e.
3.3. Crystallographic texture It is well known that FSP of face-centered cubic (FCC) metals exhibit a predominant simple shear deformation texture [30,31]. Fig. 9a shows the ideal simple shear deformation texture of FCC metals on a {111} pole figure and position of ideal orientations are also labeled. For comparative purposes, {111} pole figure of Al5083CNTs, Al5083-SiC, Al5083-SiC-CNTs, and FSPed Al5083 alloy are shown in Fig. 9cef respectively. The discrepancies between the ideal and experimental texture components are mainly due to 3 tool tilt angle and threads on tool probe which may cause a slight variation in shear plane and direction. The EBSD map (Fig. 4a) and pole figure (Fig. 9b) of base Al5083 alloy clearly shows an elongated grain structure and strong rolling texture dominated by the S components and with the minor presence of the B component with a peak intensity of 5.75. As viewed in Fig. 9cef, strong {112} <110> B/B texture component for FSPed Al5083-SiC, {111} <112> A*1/A*2 texture and {112} <110> B/B texture component for FSPed Al5083CNTs, and {001} <110> C texture, {111} <112> A*1/A*2 texture for FSPed Al5083-SiC/CNTs and {001} <110> C texture, {111} <112>A*2 texture and {111} <110> A texture components for FSPed Al5083 alloy were observed in the stir zone respectively. Addition of the reinforcements and re-precipitation of second phase particles weakened the texture due to activation of PSN mechanism and Zener pinning. Furthermore, other ideal shear component are also present to some extent in all FSPed samples. The variation in texture components in the FSPed samples is mainly due to the presence of reinforcements and their distributional effect, because the process parameters is the same for all FSPed samples. It is reported that at the lower strain value, A/A texture component is
dominant and rapidly switched to C texture component as the strain value reached between 3 and 5 and beyond which B/B texture component slowly replaces the C texture component [32,33]. 3.4. Microhardness Microhardness profiles of the FSPed regions of composites along the cross-section after three FSP passes are shown in Fig. 10. The average hardness of the base Al5083 alloy was found to be 79.25 ± 3.96 HV. It can be observed that the average hardness of FSPed Al5083 alloy was about 94.00 ± 4.70 HV, which is higher when compared with base Al5083 alloy. The factors which are responsible for the improved hardness of FSPed Al5083 alloy are grain refinement and re-precipitation of second phase particles. Initially, base Al5083 alloy has elongated grains with an average size of 49 mm (Fig. 4a) and large complex shaped second phase particles which are then replaced by fine equiaxed grain structure, fragmentation, and re-distribution of precipitates throughout the stir zone by the thermo-mechanical process of FSP. However, these second phase particles do not dissolve in the stir zone. Thus only fragmentation of particles occurs [28,29]. The aluminum 5083 alloys reinforced with SiC and CNTs either individually or in aggregate form shows an improvement in the hardness compared to FSPed Al5083 alloy and base Al5083 alloy. The average hardness of Al5083-SiC, Al5083-SiC/CNTs and Al5083-CNTs are 124.50 ± 6.22 HV, 112.50 ± 5.62 HV, and 107.50 ± 5.37 HV respectively, which is 1.5, 1.4 and 1.35 times, higher than base Al5083 alloy. The enhanced hardness is mainly due to a homogeneous distribution of CNTs and SiC particle results in dislocations pinning and impeding grain growth, good interfacial bonding between the matrix and particles. 3.5. Tensile properties Fig. 11 depicts the engineering stress-strain curves for Al5083 alloy and FSPed samples. Tensile tests were carried out to determine the tensile strength, yield strength, and elongation of FSPed samples and base Al5083 alloy. It is evident that all the FSPed composite samples exhibited higher tensile strength with a subsequent drop in the elongation as compared to base Al5083 alloy. Among the FSPed samples, SiC/CNTs reinforced Al5083 matrix composites shows the maximum tensile strength of 361 MPa. Whereas Al5083-SiC, Al5083-CNTs, FSPed Al5083, and Al5083 alloy shows the ultimate tensile strength of 335 MPa, 316 MPa, 306 MPa, and 298 MPa respectively. However, FSPed Al5083 alloy shows higher elongation compared to other FSPed samples and base
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Fig. 8. Bright field TEM images of Al5083-SiC/CNTs showing (a) Fine second phase particles in stir zone (b) dislocation rearrangement into subgrain boundaries (white arrows) (c) second phase fine particle obstruct dislocation motion, (d) survived twisted CNTs with Al4C3 phase (e) interface between CNTs and Al matrix leading to Al4C3 phase formation.
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Fig. 9. {111} pole figure of (a) ideal orientations of FCC metals under simple shear deformation texture with fiber component labeled [30,31], (b) Al5083 alloy, (c) FSPed Al5083-SiC, (d) FSPed Al5083-CNTs, (e) FSPed Al5083-SiC/CNTs and (f) FSPed Al5083 alloy.
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Hosseini et al. [18] is having good agreement with our results. Similar results of Al5083-SiC composites were reported by Amra et al. [17]. However, in this work, Al5083-CNTs/SiC composites show a higher ultimate tensile strength of 361 MPa which clearly shows that FSP is effectively utilized to fabricate AMMCs without any detrimental products. The properties difference mainly depends upon the chemical composition, processing parameters, and morphology of reinforcement used. 3.6. Fractography
Fig. 10. Microhardness profile across the transverse section of the Al5083 alloy and FSPed samples.
Fig. 12 depicts the SEM fracture surface of Al5083 alloy and FSPed samples. Fig. 12a shows the presence of larger dimples and voids indicating a typical ductile mode of failure. In the FSPed Al5083 alloy (Fig. 12b), a similar type of ductile failure with marginally larger dimples and voids are observed which indicates that FSPed Al5083 alloy exhibit higher elongation before fracture as compared to base Al5083 alloy and other FSPed composite samples. The improved strength and elongation of FSPed Al5083 alloy are attributed to grain refinement owing to dynamic recrystallization [17]. Some of the second phase intermetallic particles are present in the core of the dimples of Al5083 alloy and FSPed Al5083 alloy. The FSPed composite samples (Fig. 12c and d) show the bimodal distribution of smaller to larger dimples and voids are attributed to grain refinement and homogenous distribution of reinforcement [18,34]. During tensile loading, voids are formed around the particles and mobility of dislocation hinder by reinforcement which further increases stress to deform the material. Failure occurs due coalescence of voids, followed by growth. The fracture sites in Al5083-SiC composite sample shows that the SiC particles found at the bottom of the dimples are shown in Fig. 12d. The fracture occurs at particle/matrix interfaces showing good cohesion between the particle and matrix resulted in sound interfacial bonding. 4. Conclusions In the present study, the effect of CNTs/SiC either separately or in aggregate form, on the microstructure, texture evolution, and mechanical properties of Al5083 composites were investigated. Based on the experimental results, the following conclusions were drawn:
Fig. 11. Stress-strain curves of FSPed samples and base Al5083 alloy.
Al5083 alloy. Al5083 matrix strengthening was obtained by grain refinement, dispersion of particles and re-precipitates, and dislocations. Uniform dispersion of reinforcement, clean interface, and strong interfacial bonding enhances the dispersion strengthening. Second, fine-grained material offers more resistance to deformation and substantial strengthening by grain refining. Third, due to the difference in coefficient of thermal expansion between reinforcement and matrix induces additional dislocations during material flow. The mobility of dislocations is hinder by the reinforcement, which further enhances the strength of materials. The tensile properties of the present work are comparable with values reported in the literature. Bauri et al. [34] have reported tensile strength of 296 MPa and 337 MPa, and % elongation of 25% and 33% for base A15083 alloy and FSPed Al5083 alloy respectively. The work of Yuvaraj et al. [35] having a close agreement with present work, i.e., tensile strength and % elongation of Al5083 alloy before and after FSP are 300 MPa, 27.1%, and 312 MPa, 30.2% respectively. Also, CNTs reinforced with Al5083 alloy reported by
1. CNTs/SiC reinforced Al5083 composites are successfully processed using multi-pass FSP with a uniform distribution of SiC and CNTs. FSP also results in significant grain refinement, fragmentation, and re-precipitation of second phase particles and partially initiated an interfacial reaction between the Al matrix and CNTs to form in-situ Al4C3 intermetallic compound. 2. The dominant texture component observed in the stir zone of FSPed samples are {112} <110> B/B for Al5083-SiC, {111} <112> A*1/A*2 texture and {112} <110> B/B texture for Al5083-CNTs, {001} <110> C texture and {111} <112> A*1/A*2 texture for Al5083-SiC/CNTs. 3. Incorporation of CNTs/SiC and second phase particles resulted in the activation of Zener-Holloman mechanism and particlestimulated nucleation (PSN) mechanism by developing randomly oriented grains. 4. The average hardness of Al5083-SiC, Al5083-SiC/CNTs and Al5083-CNTs are 124.50 ± 6.22 HV, 112.50 ± 5.62 HV, and 107.50 ± 5.37 HV respectively, which is 1.5, 1.4 and 1.35 times, higher hardness than base Al5083 alloy. 5. The maximum tensile strength of 361 MPa was obtained for Al5083-SiC/CNTs hybrid composite whereas, Al5083-SiC, Al5083-CNTs, FSPed Al5083 alloy and base Al5083 alloy shows
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Fig. 12. Fracture surface (a) Al5083, (b) FSPed Al5083 alloy, (c) Al5083-SiC/CNTs (d) Al5083-CNTs, (e) Al5083-SiC.
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the tensile strength of 335 MPa, 316 MPa, 306 MPa, and 298 MPa, respectively. [17]
Acknowledgments The author (Vikram Kumar S. Jain), acknowledge Department of Science & Technology, Govt. of India for sponsoring him to pursue Ph.D. under INSPIRE Fellowship (DST/INSPIRE Fellowship/2015/ IF150488). The authors acknowledge the use of the National Facility of Texture & OIM, IIT Bombay for their support in carrying out the EBSD studies. The authors thank Dr. Devinder Yadav, Assistant Professor, IIT Patna for a useful discussion of the manuscript.
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