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Role of powder preparation route on microstructure and mechanical properties of Al-TiB2 composites fabricated by accumulative roll bonding (ARB)
Materials Science & Engineering A 677 (2016) 400–410
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Role of powder preparation route on microstructure and mechanical properties of Al-TiB2 composites fabricated by accumulative roll bonding (ARB) M. Askarpour, Z. Sadeghian n, M. Reihanian Department of Materials Science and Engineering, Faculty of Engineering, Shahid Chamran University of Ahvaz, Ahvaz, Iran
art ic l e i nf o
a b s t r a c t
Article history: Received 13 July 2016 Received in revised form 5 September 2016 Accepted 17 September 2016 Available online 18 September 2016
Accumulative roll bonding (ARB) was conducted up to seven cycles to fabricate Al-TiB2 particulate metal matrix composites. The reinforcing particles were prepared and used in three different processing conditions: as-received TiB2, mixed TiB2-Al and in-situ synthesized TiB2-Al. The mixed TiB2-Al powder was produced by milling of TiB2 with Al powder and in-situ synthesized TiB2-Al powder was prepared by mechanical alloying (MA) through inducing TiB2 particles in the Al with various composition of 10, 20 and 30 wt% Al. Transmission electron microscope (TEM) and scanning electron microscope (SEM) were used to evaluate the microstructure of the produced composites. The composite obtained from the insitu TiB2-Al powder showed the most uniform distribution of particles and exhibited the highest tensile strength of about 177 MPa in comparison with the composites reinforced with the as-received TiB2 (156 MPa) and mixed TiB2-Al powder (160 MPa). After seven ARB cycles, an ultra-fine grained structure with the average size of about 300 nm was obtained in the composite reinforced with in-situ TiB2-Al powder. The appearance of dimples in tensile fracture surfaces revealed a ductile-type fracture in the produced composites. & 2016 Elsevier B.V. All rights reserved.
Keywords: Accumulative roll bonding (ARB): Metal matrix composite Mechanical alloying (MA) Ultra-fine grained structure
1. Introduction In recent decades, Al-based metal matrix composites (MMCs) have received increasing attention in a wide range of applications due to their attractive properties, such as high specific strength, stiffness, and wear resistance compared with the unreinforced alloys [1]. Several techniques consisting powder metallurgy and casting have been developed to fabricate particle reinforced MMCs [2,3]. It has been demonstrated that manufacturing method can affect the interfacial particle-matrix bond strength and mechanical properties of the composite [4]. Recently, the accumulative roll bonding (ARB) process has been developed as a severe plastic deformation (SPD) technique to fabricate ultra-fine grained metals [5]. More recently, this method has been used metal to manufacture particulate MMCs [6] because of several advantageous compared with the conventional fabrication techniques. For example, it is possible to obtain ultra-fine grained composites in the form of high strength sheets with a faster production rate. Meanwhile, protective atmosphere is not necessary and the process is performed at room temperature in comparison to powder n
Corresponding author. E-mail address: [email protected] (Z. Sadeghian).
http://dx.doi.org/10.1016/j.msea.2016.09.068 0921-5093/& 2016 Elsevier B.V. All rights reserved.
metallurgy methods. ARB has been successfully employed for fabrication of MMCs with different conventional reinforcing ceramic particles such as SiC [7,8], Al2O3 [9–11], B4C [12] and Fe3O4 [13] and hybrid composites containing two reinforcing particles [14,15]. It has been shown that a uniform distribution of particles can be obtained after imposing a critical strain during ARB [16,17]. Though inducing ceramic particles results in the improvement of the strength and hardness, they mostly exhibit several disadvantages such as poor bonding with the matrix and the tendency to agglomerate and to form clusters. TiB2 as a refractory compound is an attractive reinforcement for a range of applications including high temperature components because it possesses a high modulus, excellent refractory properties and chemical inertness [18]. TiB2 is particularly suitable as a reinforcing phase for Al matrix composites because of its thermodynamic stability in Al matrix [19]. There are a few routes to synthesize Al-TiB2 composites, but in situ approach is particularly very attractive. In situ composites are multiphase materials where the reinforcing phase is synthesized within the matrix during the composite fabrication. The clean interface resulting from the absence of oxidation, during in situ formation of reinforcement, offers the potential to improve the strength. The strong bonding between the reinforcing particles and the Al matrix has been verified to affects the mechanical
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properties of the composite [20]. The in-situ synthesized reinforcing particles are finer in size and have a uniform distribution. In addition, Al matrix composite containing the in-situ TiB2 reinforcing particles can be an appropriate choice for high strength applications. The matrix-reinforcement interface in the in-situ produced TiB2 is reported to be parallel to TiB2 lattice base plane [21]. Over the last decades, the in-situ methods have been rigorously investigated to fabricate MMCs with enhanced properties [22]. One method of in-situ processing is mechanical alloying (MA) which is considered as a solid-state powder processing method for fabricating the MMCs. The advantage of this process is the production of nano structured composites with appropriate distribution of reinforcements in the matrix [23]. In situ fabrication of TiB2 reinforced MMCs by MA process has been investigated in a few studies [24–26]. According to the thermodynamic data, TiB2 is the most stable phase in comparison to TiAl and Al3Ti intermetallic compounds [27]. It has been reported that Al content can affect the kinetic of reaction during MA, resulting in the formation of unfavorable phases such as Al3Ti [24]. The effect of grain boundary complexions on material characteristics like sintering behavior or grain growth has been investigated in a few metal-ceramic systems [28]. The present paper is an effort to compare the effect of in-situ synthesized and ex-situ TiB2 reinforcing particles on mechanical properties of Al/TiB2 composites produced by ARB. The effect of the powder preparation was investigated by conducting three different conditions: as-received TiB2, mixed TiB2-Al and in-situ synthesized TiB2-Al. The insitu synthesized TiB2-Al powder was prepared by 20 h MA with different compositions of Al (10, 20 and 30 wt%) and the optimum composition was used for ARB processing. The microstructure and mechanical properties of the composites obtained from different processing routes were examined and compared with each other.
2. Materials and method TiB2 (99.9,o 10 μm) powder was provided as the ex-situ reinforcing particle. As-received TiB2 powder was mixed and milled with Al powder for 5 h to obtain TiB2-20 wt% Al composite powder. Elemental powders of Al (99.8%, 63 μm), Ti (99.9%, 40–60 μm) and B (4 99%, 2 μm) were used as the starting materials for the insitu fabrication of TiB2. Mixtures of 61.97Ti–28.03B–10Al, 55.09Ti24.91B–20Al and 48.2Ti-21.8B–30Al were milled for 20 h to achieve in situ TiB2-Al composite powders by MA. Powder mixtures were milled by a Fritsch type planetary ball mill with a rotating speed of 350 rpm. The ball to powder weight ratio was chosen to be 10:1 and the diameter of the chromium steel balls was 15 mm. The hardened chromium steel vial was evacuated and filled with pure argon at atmospheric pressure to prevent oxidation during the MA process. In order to avoid severe adhesion of Al powder to the balls and the vial surfaces, 1 wt% zinc stearate was added to the mixture as a process control agent. As-received commercially pure Al (AA1050) in the form of sheets with 1 mm thickness was used as the matrix material. For ARB processing, the strips were cut into 150 mm ×50 mm and annealed at 400 C for 2 h. To set the prepared surfaces in contact and closely fixed to each other, four holes were drilled near the edges of the strips. The annealed strips were degreased by acetone and scratch brushed with a circular stainless steel brush having a 0.3 mm wire diameter. To fabricate the Al-TiB2 composites, ARB process was carried out in two steps. The first step includes three cycles and the second step contains four cycles. In the first cycle, 0.5 vol% of TiB2 powders was uniformly distributed between three Al strips (about 0.25 vol% between each layer). The stacked strips were fastened at both ends by copper wires and rolled through a 60% reduction in thickness at ambient temperature without
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Fig. 1. XRD patterns of different Ti-B-Al powder mixtures, a) TiB2-20 wt%Al (mixed), b) TiB2 30 wt%Al (in-situ), c) TiB2 20 wt%Al (in situ) and d) TiB2 10 wt%Al (in situ).
lubrication. In the second and the third cycle, the roll-bonded strip was cut into two halves while 0.25 vol% of TiB2 powders dispersed between them. The stacked strips were roll-bonded through a 50% reduction in thickness. After this step, the total amount of the TiB2 reached to about 2% vol% and the thickness of the composite sheet was about 1 mm. In the second step, ARB was repeated up to four cycles (seven cycles in total) without adding the particles between the layers. ARB was conducted with a rolling machine having the capacity of 30 tons and rolling diameter of 170 mm. The rolling speed was set to 4 rpm. For simplicity the composites produced by the as-received, mixed Al-TiB2, and in-situ Al-TiB2 powders are referred to Al-TiB2 (as-received), Al-TiB2 (mixed) and Al-TiB2 (insitu) composite, respectively. A SEIFERT 30033 PTS diffractometer employing monochromatic Cu Kα1 radiation (λ ¼ 0.15406 nm) was used to investigate structural changes after MA. XRD scans were performed with a step size of 0.05° in 2θ and a dwell time per step of 20 s. The MIRA3 TESCAN field emission scanning electron microscope (SEM) equipped with energy dispersive spectroscopy (EDS) was utilized for investigating the microstructure on rolling direction-normal direction (RD-ND) plane and rolling direction-transverse direction (RD-TD) plane and the cross-sectional microstructure of the powder particles. The powders were mounted within a conductive mounting resin and the mounted samples were prepared by metallography techniques for microstructural investigation. Transmission electron microscopy (TEM) specimens were cut from the ARB processed samples by a micro-cutter. The surface of the specimens was mechanically polished to the thickness of 100 μm using SiC abrasive papers. After mechanical polishing, thin foils were made by twin-jet polishing Tenupole 5 facility (Struers Co., Ltd.) in a 200 ml HClO4 and 800 ml ethanol solution with the applied voltage of 35 V at 20 C for perforation. The specimens were then polished by ion beam using Gatan 691 precision ion polishing system (PIPS). For TEM/STEM, a transmission electron microscope (STEM, JEOL JEM-2100F) with the acceleration voltage of 200 kV was used. Tensile samples were machined from the ARB processed strips, according to the ISO 6892-1 standard, with orientation along the rolling direction. The gage width and length of the samples were 5 and 10 mm, respectively. Tensile samples were tested at room temperature with a tensile testing machine (SANTAM STM-150) under a nominal initial strain rate of 10 3 s 1.
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Fig. 2. Cross sectional BSE micrographs and EDS point analysis of a) TiB2 30 wt%Al, b) TiB2 20 wt%Al and TiB2 10 wt%Al powder mixtures obtained from 20 h of MA.
3. Results and discussion 3.1. Powder preparation Fig. 1 shows the X-ray diffraction patterns of Al-Ti-B powder mixtures with different Al contents after 20 h of MA. XRD pattern of TiB2 20 wt%Al (mixed) is also presented that shows the stability of as-received TiB2 after 5 h of milling with Al. Peaks related to TiB2 appears on the XRD pattern of MAed TiB2 30 wt%Al powder. However, the existence of Ti peaks shows the incomplete
reaction between Ti and B with increasing the Al content. The absence of boron peaks is due to its amorphous structure. Ti peaks disappear in powder mixtures containing 10 and 20 wt% Al, suggesting the completion of the Ti-B solid state reaction. The main reasons for disappearing the Al peaks in the XRD pattern of TiB2 20 wt%Al and TiB2 10 wt%Al powders are the high strain level induced in the Al lattice, nano-sized Al grains and lower X-ray scattering intensity of Al [23]. Structural evolutions and mechanism of TiB2 formation by MA process have been discussed in details, elsewhere [29].
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Fig. 3. Secondary electron micrographs of initial powders for ARB processing; a) TiB2 (as-received), b) Al- TiB2 (mixed) and c) Al-TiB2 (in situ).
Fig. 2 shows the cross-sectional SEM micrographs of typical powder particles obtained from different powder mixtures after 20 h of MA. Microstructure of TiB2 30 wt%Al shows a distribution of light and dark areas in Al matrix (Fig. 2a). EDS analysis at the exhibited points A and B reveals that the small dark areas in the matrix are B-rich, while the white fragments are Ti-rich remnants. According to the thermodynamic data, TiB2 is the most stable phase in Al-Ti-B system. Nevertheless, it seems that the presence of a large fraction of Al can limit the inter-diffusion of Ti-B in the powder mixture; preventing the completion of TiB2 synthesis. Suppressive effect of large Al amount on Ti-B reaction in Al-Ti-B system has been reported previously [24]. Typical cross-sectional SEM micrograph of TiB2 20 wt%Al particles after a total milling time of 20 h shows the formation of fine particles which are, according to EDS point analysis at the depicted point C, TiB2 particles (Fig. 2b). It can be concluded that the diffusion path between Ti and B decreases with decreasing the Al content, which can assist the reaction between them and formation of TiB2. SEM image of MAed TiB2 10 wt%Al particles (Fig. 2c) indicates that it consists of TiB2 particles with a lower amount of Al as the binding element. Moreover, the size of TiB2 particles seems to be larger in this powder compared to that in TiB2 20 wt%Al composite powder. Since the Al binder can facilitate the distribution of TiB2 particles during ARB, the TiB2 20 wt%Al composite powder obtained from 20 h of MA seems to be the appropriate powder for ARB. The Ti-B reaction is completed in this powder mixture with the highest possible Al
content. Fig. 3 shows the morphology of powders utilized for ARB process. The size range of as-received TiB2 powder is roughly less than 10 μm with mostly angular particles (Fig. 3a). After milling with 20 wt% Al powder for 5 h, a mixture of fine angular particles is achieved with a size of less than 2 μm (Fig. 3b). The in-situ TiB2 20 wt%Al powder obtained by 20 h of MA shows nearly spherical submicron particles with many particles having the size less than 500 nm (Fig. 3c). 3.2. ARB processing 3.2.1. Microstructural evolutions To investigate the bonding development between layers and the distribution of particles during ARB, the microstructure of AlTiB2 (as-received) composite after various ARB cycles is selected as typical and the results are presented in Fig. 4. After the first and the third cycle (Fig. 4a and b), the particles are seen at layer interfaces and a non-uniform distribution of large clusters of TiB2 particles as well as particle free zones can be observed. After five ARB cycles (Fig. 4c), particles are removed from interfaces to some extent, but the distribution of TiB2 particles in Al matrix seems still non-uniform. As the ARB is continued up to seven cycles (Fig. 4d), the interface between Al layers cannot be easily detected and a nearly uniform distribution of TiB2 particles in Al matrix is obvious. In this stage, the particles are separated from layer interfaces and moved into the matrix. This can be explained by film
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Fig. 4. Low magnification SEM micrographs of the Al-TiB2 (as-received) composite after various ARB cycles; (a) one, (b) three, (c) five and (d) seven cycles.
theory as follows [30]. During ARB process, two opposing work hardened layers or brittle surface oxide layers on Al strips break up coherently and the underlying metals extrude through the produced channels between the cracks from both sides of the interfaces. The presence of particles at interfaces can be considered as a preventing barrier for the film theory [31]. However, the Al matrix deforms plastically through the particle clusters and fresh metal surfaces can be obtained during ARB. At the same time, the rolling pressure can generate sufficient force and good bonding at layer interfaces even in the presence of ceramic particles. At the same time, with increasing the ARB cycles, the number of Al layers increases while their thickness decreases, leading to the further distribution of the particles and their movement from the interfaces into the bulk of the Al matrix. It is noted that the low magnification SEM image of the composite (as presented in Fig. 4) is not appropriate to show the bonding quality of the individual particles with the Al matrix. It just explains the fragmentation of the particles at interfaces and their distribution in the Al matrix. High magnification SEM micrographs of Al-TiB2 (in-situ) composite are presented in Fig. 5 to indicate the structure of TiB2 fragments and the quality of bonding between the particle and matrix after different ARB cycles. After the first ARB cycle, the material is severely porous at the interfaces with a weak bonding between the particle and matrix. These pores may be between either in the layer interface or near the Al and the particle interfaces. With increasing the number of cycles, the amount and size of porosities decreases and bonding at interfaces improves significantly. Fig. 6 presents the SEM micrographs of various Al-TiB2 composites on RD-TD plane after three and seven ARB cycles. In all composites, clusters of TiB2 particles as well as particle free zones can be seen apparently after three ARB cycles. The agglomeration of particles is reduced to some extent after seven ARB cycles. It has
been stated that the distance between the particles inside the clusters can increase due to the elongation of the clusters in rolling direction [8,32]. Consequently, the dense clusters are changed into diffuse ones with islands of Al matrix between them [33]. The microstructure of Al-TiB2 composites in the as-received and premixed conditions reveals that by premixing of TiB2 particles with Al, improved distribution with less agglomeration of TiB2 particles can be achieved after each ARB cycle. It can be concluded that the existence of Al binder between TiB2 particles in the mixed Al-TiB2 powder can facilitate the extrusion of Al matrix through initial clusters of particles. The composite with the in situ Al-TiB2 powder exhibits the best distribution as well as the least agglomeration among the others after the third and the seventh cycle. It has been reported that the in-situ synthesis results in clean and defect free interfaces between the in-situ reinforcing particles and the adjacent phase. This can provide a strong interfacial bonding between the particles and the matrix and can improve the mechanical properties of the composite [22]. TEM micrographs and the corresponding selected area diffraction (SAD) patterns (take with 1.3 mm aperture size) on RD-TD plane of Al-TiB2 (in situ) composite are shown in Fig. 7. After three ARB cycles, the grains are divided into subgrains through formation of the dislocation walls. SAD pattern of the sample shows a near net pattern, indicating that the boundaries have low misorientation in the corresponding area. Contours or fringes that can be seen mostly near the boundaries reflect the internal stress existing in grains. After seven ARB cycles (Fig. 7b), a significant decrease in grain size is observed and the microstructure consists of an array of ultrafine equiaxed grains with a less regular arrangement. SAD pattern shows the ring like patterns consisting of separate spots that indicates there are higher fractions of boundaries with high angles of misorientation. However, the size of grains seems to vary in different areas that can be attributed to the
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Fig. 5. SEM micrographs of Al-TiB2 (in situ) composite after various ARB cycles; (a) three, (b) five, (c) seven cycles and d) EDS analysis of point A.
distribution of TiB2 particles within the matrix. Generally, the formation of the ultra-fine structure in the highly strained materials is regionally inhomogeneous [34] because of the friction between the rolls and the strip surfaces. The sequence of grain refinement during SPD process has been explored by several researchers as follows [35]. A fine dislocation cell structures with low-angle grain boundaries is created by arrangement of dislocations into the low energy dislocation structure [36]. The dislocation cell structure or subgrains become finer and their misorientation angle increases due to operation of multi-directional slip and absorption of further dislocations [37]. Finally, a distinct grain structure with distinguishable high-angle boundaries is generated when the sample is subjected to a very intense plastic strain. Fig. 8 shows the STEM micrographs of Al-TiB2 (in situ) composite after three and seven ARB cycles. After three cycles of ARB, the structure contains grains with average size of about 850 nm and includes the subgrains with dislocation cell structures. The dislocations can be observed in grains with different geometries. Some of them accumulate as tangles of multiple dislocations and some emerge as individual dislocations in coarse grains (indicated in Fig. 8a). The contrast change in some grains can be considered as an evidence for the initiation of subgrain formation. The presence of subgrains within the grains is an indication for mechanism of continuous dynamic recrystallization mechanism [38]. This mechanism is based on continuous absorption of dislocations in subgrain boundaries, followed by the formation of new grains separated by high angle boundaries. After seven ARB cycles, the average grain size is decreased to about 300 nm. The dislocation density inside the grains is very low in comparison with that
observed after three ARB cycles. It has been stated that the development of fine structures during SPD is accompanied by a decrease in dislocation density at large strains. Indeed, severely deformed materials commonly contain fine grains that are completely free of dislocations [35]. Extinction of dislocations inside the grains in the final stages of ARB has been reported by previous researcher [39]. 3.2.2. Mechanical properties Fig. 9a illustrates the representative engineering stress–strain curves of the annealed Al and Al-TiB2 (in situ) composite after different cycles of ARB processing. In general, the tensile test is not an appropriate tool to analyze the elastic properties of the materials and more accurate test methods are required. Therefore, the engineering stress-strain curves obtained from tensile test do not exhibit a significant difference between the elastic modulus of the composite materials produced by ARB. A similar trend can also be observed in literature such as in Al/SiC [7], Al/CuP [40] and Al/ Al2O3/SiC composite [41]. The variation in the ultimate tensile strength (UTS) and uniform elongation with the number of cycles is also plotted in Fig. 9b. After the first cycle, a significant increase in UTS to about 155 MPa is observed in comparison to that of the annealed Al (47 MPa). The UTS increases gradually with increasing the number of ARB cycles up to 171 MPa after seven ARB cycles. Tensile curves of the ARB processed composite show immediate UTS after yielding unlike the annealed sample. Contrary to the strength, the elongation decreases drastically after the first ARB cycle. During tension test, the response of nano/ultra-fine grained polycrystalline metals is different from that of coarse grain metals.
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Fig. 6. SEM micrographs of various composites; a), c), e) three cycles and b), d), f) seven cycles.
Several reports show that the work hardening is absent in ultrafine grained materials that are produced by severe plastic deformation [42]. This leads virtually to necking at the yield stress (in tension), and the net results is a low tensile elongation. In the other words, the yield stress and UTS of SPD processed materials are almost identical. Therefore, the variation of UTS with the number of ARB cycle (that is presented in Fig. 9b.) can be considered as a good approach to show the effect of ARB on the yield stress. Two important mechanisms that are responsible for the deformation behavior of the SPD processed materials are strain hardening and grain refinement. The former is dominated in the
early stages and the latter is occurred in the last stages of SPD [43]. The increase in dislocation density within the grains and formation of low angle dislocation boundaries are the main reasons of strain hardening in the early stages of ARB. However, in the final stages, the size of dislocation boundaries decreases, the misorientation at these boundaries increases progressively and grain refinement to the ultra-fine scale causes the increase in the strength [44]. In addition, the existence of TiB2 reinforcing particles can also be responsible as additional barriers for moving dislocations and consequently increase in the tensile strength. Considering the Orowan strengthening mechanism that involves dislocations bowing around the particles, dispersion of fine
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(a)
(b)
Fig. 7. TEM micrographs and corresponding SAD patterns of Al-TiB2(in situ) composite after a) three and b) seven ARB cycles.
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Fig. 8. STEM micrographs of Al-TiB2(in situ) composite after a) three and b) seven ARB cycles.
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Fig. 10. The engineering stress–strain curves of final Al-TiB2 composites obtained from seven ARB cycles.
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particles is an effective factor for the strengthening. The extreme decrease in the elongation can be explained by the increase in density of dislocation at the first ARB stages that promotes the nucleation of cracks. Moreover, TiB2 reinforcing particles can induce voids and cracks that can reduce the ductility of the composite. However, a slight increase in elongation after successive ARB cycles can be seen which can be resulted from enhancement
of the bonding at interfaces. In order to compare the deformation behavior of Al-TiB2 composites, the engineering stress–strain curves of all composites (after seven ARB cycles) are presented in Fig. 10. The Al-TiB2 (in situ) composite shows the largest UTS (171 MPa) while the least UTS belongs to Al-TiB2 (as-received) (156 MPa). Different characteristics of the reinforcing particles such as morphology, size, quantity and distribution can affect the strength of the composites. Since the Al-TiB2 (in situ) composite reveals a more uniform distribution of fine TiB2 particles with less porosity and spherical morphology, the highest UTS is expected for this composite. Defects such as porosity at particle-matrix interface and agglomeration of particles are considered as main reasons for composite failure. Fig. 11 presents typical fracture surfaces of Al-TiB2 (in situ) composite after three and seven ARB cycles. The fracture surfaces show a typical ductile fracture that is characterized by dimples. This mode of fracture occurs by the formation and coalescence of microvoids, crack propagation and final fracture [45]. Several voids can be elongated in one direction because of unequal triaxial stresses [46]. The presence of particles at the bottom of some dimples, indicated in the micrographs, can be considered as an evidence for the nucleation of voids from particle-matrix interfaces [47]. The agglomeration of particles after the third cycle
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Fig. 11. SEM micrographs of the fracture surfaces after tensile experiment of Al-TiB2 (in situ) composite obtained from (a) three and (b) seven ARB cycles.
(Fig. 11a) is obvious on the fracture surface while the no particle clustering can be seen after the seventh cycle (Fig. 11b). The presence of the particles can change the morphology of dimples in a manner that the dimples nucleated at the particle sites (as indicated in Fig. 11b) are deeper and larger than that nucleated in other regions. These results are consistent with those obtained in the previous sections (Section 3.2.2) where a uniform distribution of particles was observed after seven ARB cycles. It is noted that the dimples are no longer apparent at regions where the particles clustering occur; indicating that agglomeration of particles can affect the mode of fracture as well as mechanical properties.
5. The fracture surface of Al-TiB2 (in situ) composite shows several dimples, which is an indication of ductile-type fracture. The agglomeration of particles on the fracture surface of the composite after the third cycle is more evident.
Acknowledgment Financial support provided by Shahid Chamran University of Ahvaz (Grant no. 94-3-02-31579) is gratefully acknowledged. Authors would like to thank Mr. E. Bagherpour for the preparation of TEM samples and taking TEM images in Metallic Materials Science Laboratory of Doshisha University, Japan.
4. Conclusions Al-TiB2 metal matrix composites were fabricated by ARB up to seven cycles. The reinforcing particles were prepared in three different processing conditions: as-received TiB2, mixed TiB2-Al and in-situ synthesized TiB2-Al. The most important conclusions are: 1. In-situ synthesized TiB2-Al powder with 10, 20 and 30 wt% Al is prepared by 20 h MA. It is concluded that the in-situ synthesized TiB2 20 wt%Al can be the appropriate powder for ARB because of completing the Ti-B solid state reaction. The higher inter-diffusion between Ti and B in the powder mixture (compared with TiB2 30 wt%Al) and the smaller size of TiB2 particles (compared with TiB2 10 wt%Al) are the advantages of TiB2 20 wt%Al powder mixture. 2. In comparison to the Al-TiB2 (as-received) and Al-TiB2 (mixed) composites, the best distribution as well as the least agglomeration is obtained in the Al-TiB2 (in-situ) composite. It was attributed to the formation of the clean, defect free interfaces between the in-situ produced powders and the adjacent phase. 3. After three ARB cycles, a subgrain structure with an average grain size of about 850 nm is obtained in the Al-TiB2 (in-situ) composite. ARB processing up to seven cycles leads to an ultrafine grained structure with an average size of about 300 nm in the Al-TiB2 (in-situ) composite. 4. The highest tensile strength of about 171 MPa is achieved for the Al-TiB2 (in situ) composite because of a more uniform distribution of fine TiB2 particles with less porosity and spherical morphology.
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Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Role of powder preparation route on microstructure and mechanical properties of Al-TiB2 composites fabricated by accumulative roll bonding (ARB) M. Askarpour, Z. Sadeghian n, M. Reihanian Department of Materials Science and Engineering, Faculty of Engineering, Shahid Chamran University of Ahvaz, Ahvaz, Iran
art ic l e i nf o
a b s t r a c t
Article history: Received 13 July 2016 Received in revised form 5 September 2016 Accepted 17 September 2016 Available online 18 September 2016
Accumulative roll bonding (ARB) was conducted up to seven cycles to fabricate Al-TiB2 particulate metal matrix composites. The reinforcing particles were prepared and used in three different processing conditions: as-received TiB2, mixed TiB2-Al and in-situ synthesized TiB2-Al. The mixed TiB2-Al powder was produced by milling of TiB2 with Al powder and in-situ synthesized TiB2-Al powder was prepared by mechanical alloying (MA) through inducing TiB2 particles in the Al with various composition of 10, 20 and 30 wt% Al. Transmission electron microscope (TEM) and scanning electron microscope (SEM) were used to evaluate the microstructure of the produced composites. The composite obtained from the insitu TiB2-Al powder showed the most uniform distribution of particles and exhibited the highest tensile strength of about 177 MPa in comparison with the composites reinforced with the as-received TiB2 (156 MPa) and mixed TiB2-Al powder (160 MPa). After seven ARB cycles, an ultra-fine grained structure with the average size of about 300 nm was obtained in the composite reinforced with in-situ TiB2-Al powder. The appearance of dimples in tensile fracture surfaces revealed a ductile-type fracture in the produced composites. & 2016 Elsevier B.V. All rights reserved.
Keywords: Accumulative roll bonding (ARB): Metal matrix composite Mechanical alloying (MA) Ultra-fine grained structure
1. Introduction In recent decades, Al-based metal matrix composites (MMCs) have received increasing attention in a wide range of applications due to their attractive properties, such as high specific strength, stiffness, and wear resistance compared with the unreinforced alloys [1]. Several techniques consisting powder metallurgy and casting have been developed to fabricate particle reinforced MMCs [2,3]. It has been demonstrated that manufacturing method can affect the interfacial particle-matrix bond strength and mechanical properties of the composite [4]. Recently, the accumulative roll bonding (ARB) process has been developed as a severe plastic deformation (SPD) technique to fabricate ultra-fine grained metals [5]. More recently, this method has been used metal to manufacture particulate MMCs [6] because of several advantageous compared with the conventional fabrication techniques. For example, it is possible to obtain ultra-fine grained composites in the form of high strength sheets with a faster production rate. Meanwhile, protective atmosphere is not necessary and the process is performed at room temperature in comparison to powder n
Corresponding author. E-mail address: [email protected] (Z. Sadeghian).
http://dx.doi.org/10.1016/j.msea.2016.09.068 0921-5093/& 2016 Elsevier B.V. All rights reserved.
metallurgy methods. ARB has been successfully employed for fabrication of MMCs with different conventional reinforcing ceramic particles such as SiC [7,8], Al2O3 [9–11], B4C [12] and Fe3O4 [13] and hybrid composites containing two reinforcing particles [14,15]. It has been shown that a uniform distribution of particles can be obtained after imposing a critical strain during ARB [16,17]. Though inducing ceramic particles results in the improvement of the strength and hardness, they mostly exhibit several disadvantages such as poor bonding with the matrix and the tendency to agglomerate and to form clusters. TiB2 as a refractory compound is an attractive reinforcement for a range of applications including high temperature components because it possesses a high modulus, excellent refractory properties and chemical inertness [18]. TiB2 is particularly suitable as a reinforcing phase for Al matrix composites because of its thermodynamic stability in Al matrix [19]. There are a few routes to synthesize Al-TiB2 composites, but in situ approach is particularly very attractive. In situ composites are multiphase materials where the reinforcing phase is synthesized within the matrix during the composite fabrication. The clean interface resulting from the absence of oxidation, during in situ formation of reinforcement, offers the potential to improve the strength. The strong bonding between the reinforcing particles and the Al matrix has been verified to affects the mechanical
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properties of the composite [20]. The in-situ synthesized reinforcing particles are finer in size and have a uniform distribution. In addition, Al matrix composite containing the in-situ TiB2 reinforcing particles can be an appropriate choice for high strength applications. The matrix-reinforcement interface in the in-situ produced TiB2 is reported to be parallel to TiB2 lattice base plane [21]. Over the last decades, the in-situ methods have been rigorously investigated to fabricate MMCs with enhanced properties [22]. One method of in-situ processing is mechanical alloying (MA) which is considered as a solid-state powder processing method for fabricating the MMCs. The advantage of this process is the production of nano structured composites with appropriate distribution of reinforcements in the matrix [23]. In situ fabrication of TiB2 reinforced MMCs by MA process has been investigated in a few studies [24–26]. According to the thermodynamic data, TiB2 is the most stable phase in comparison to TiAl and Al3Ti intermetallic compounds [27]. It has been reported that Al content can affect the kinetic of reaction during MA, resulting in the formation of unfavorable phases such as Al3Ti [24]. The effect of grain boundary complexions on material characteristics like sintering behavior or grain growth has been investigated in a few metal-ceramic systems [28]. The present paper is an effort to compare the effect of in-situ synthesized and ex-situ TiB2 reinforcing particles on mechanical properties of Al/TiB2 composites produced by ARB. The effect of the powder preparation was investigated by conducting three different conditions: as-received TiB2, mixed TiB2-Al and in-situ synthesized TiB2-Al. The insitu synthesized TiB2-Al powder was prepared by 20 h MA with different compositions of Al (10, 20 and 30 wt%) and the optimum composition was used for ARB processing. The microstructure and mechanical properties of the composites obtained from different processing routes were examined and compared with each other.
2. Materials and method TiB2 (99.9,o 10 μm) powder was provided as the ex-situ reinforcing particle. As-received TiB2 powder was mixed and milled with Al powder for 5 h to obtain TiB2-20 wt% Al composite powder. Elemental powders of Al (99.8%, 63 μm), Ti (99.9%, 40–60 μm) and B (4 99%, 2 μm) were used as the starting materials for the insitu fabrication of TiB2. Mixtures of 61.97Ti–28.03B–10Al, 55.09Ti24.91B–20Al and 48.2Ti-21.8B–30Al were milled for 20 h to achieve in situ TiB2-Al composite powders by MA. Powder mixtures were milled by a Fritsch type planetary ball mill with a rotating speed of 350 rpm. The ball to powder weight ratio was chosen to be 10:1 and the diameter of the chromium steel balls was 15 mm. The hardened chromium steel vial was evacuated and filled with pure argon at atmospheric pressure to prevent oxidation during the MA process. In order to avoid severe adhesion of Al powder to the balls and the vial surfaces, 1 wt% zinc stearate was added to the mixture as a process control agent. As-received commercially pure Al (AA1050) in the form of sheets with 1 mm thickness was used as the matrix material. For ARB processing, the strips were cut into 150 mm ×50 mm and annealed at 400 C for 2 h. To set the prepared surfaces in contact and closely fixed to each other, four holes were drilled near the edges of the strips. The annealed strips were degreased by acetone and scratch brushed with a circular stainless steel brush having a 0.3 mm wire diameter. To fabricate the Al-TiB2 composites, ARB process was carried out in two steps. The first step includes three cycles and the second step contains four cycles. In the first cycle, 0.5 vol% of TiB2 powders was uniformly distributed between three Al strips (about 0.25 vol% between each layer). The stacked strips were fastened at both ends by copper wires and rolled through a 60% reduction in thickness at ambient temperature without
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Fig. 1. XRD patterns of different Ti-B-Al powder mixtures, a) TiB2-20 wt%Al (mixed), b) TiB2 30 wt%Al (in-situ), c) TiB2 20 wt%Al (in situ) and d) TiB2 10 wt%Al (in situ).
lubrication. In the second and the third cycle, the roll-bonded strip was cut into two halves while 0.25 vol% of TiB2 powders dispersed between them. The stacked strips were roll-bonded through a 50% reduction in thickness. After this step, the total amount of the TiB2 reached to about 2% vol% and the thickness of the composite sheet was about 1 mm. In the second step, ARB was repeated up to four cycles (seven cycles in total) without adding the particles between the layers. ARB was conducted with a rolling machine having the capacity of 30 tons and rolling diameter of 170 mm. The rolling speed was set to 4 rpm. For simplicity the composites produced by the as-received, mixed Al-TiB2, and in-situ Al-TiB2 powders are referred to Al-TiB2 (as-received), Al-TiB2 (mixed) and Al-TiB2 (insitu) composite, respectively. A SEIFERT 30033 PTS diffractometer employing monochromatic Cu Kα1 radiation (λ ¼ 0.15406 nm) was used to investigate structural changes after MA. XRD scans were performed with a step size of 0.05° in 2θ and a dwell time per step of 20 s. The MIRA3 TESCAN field emission scanning electron microscope (SEM) equipped with energy dispersive spectroscopy (EDS) was utilized for investigating the microstructure on rolling direction-normal direction (RD-ND) plane and rolling direction-transverse direction (RD-TD) plane and the cross-sectional microstructure of the powder particles. The powders were mounted within a conductive mounting resin and the mounted samples were prepared by metallography techniques for microstructural investigation. Transmission electron microscopy (TEM) specimens were cut from the ARB processed samples by a micro-cutter. The surface of the specimens was mechanically polished to the thickness of 100 μm using SiC abrasive papers. After mechanical polishing, thin foils were made by twin-jet polishing Tenupole 5 facility (Struers Co., Ltd.) in a 200 ml HClO4 and 800 ml ethanol solution with the applied voltage of 35 V at 20 C for perforation. The specimens were then polished by ion beam using Gatan 691 precision ion polishing system (PIPS). For TEM/STEM, a transmission electron microscope (STEM, JEOL JEM-2100F) with the acceleration voltage of 200 kV was used. Tensile samples were machined from the ARB processed strips, according to the ISO 6892-1 standard, with orientation along the rolling direction. The gage width and length of the samples were 5 and 10 mm, respectively. Tensile samples were tested at room temperature with a tensile testing machine (SANTAM STM-150) under a nominal initial strain rate of 10 3 s 1.
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Fig. 2. Cross sectional BSE micrographs and EDS point analysis of a) TiB2 30 wt%Al, b) TiB2 20 wt%Al and TiB2 10 wt%Al powder mixtures obtained from 20 h of MA.
3. Results and discussion 3.1. Powder preparation Fig. 1 shows the X-ray diffraction patterns of Al-Ti-B powder mixtures with different Al contents after 20 h of MA. XRD pattern of TiB2 20 wt%Al (mixed) is also presented that shows the stability of as-received TiB2 after 5 h of milling with Al. Peaks related to TiB2 appears on the XRD pattern of MAed TiB2 30 wt%Al powder. However, the existence of Ti peaks shows the incomplete
reaction between Ti and B with increasing the Al content. The absence of boron peaks is due to its amorphous structure. Ti peaks disappear in powder mixtures containing 10 and 20 wt% Al, suggesting the completion of the Ti-B solid state reaction. The main reasons for disappearing the Al peaks in the XRD pattern of TiB2 20 wt%Al and TiB2 10 wt%Al powders are the high strain level induced in the Al lattice, nano-sized Al grains and lower X-ray scattering intensity of Al [23]. Structural evolutions and mechanism of TiB2 formation by MA process have been discussed in details, elsewhere [29].
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Fig. 3. Secondary electron micrographs of initial powders for ARB processing; a) TiB2 (as-received), b) Al- TiB2 (mixed) and c) Al-TiB2 (in situ).
Fig. 2 shows the cross-sectional SEM micrographs of typical powder particles obtained from different powder mixtures after 20 h of MA. Microstructure of TiB2 30 wt%Al shows a distribution of light and dark areas in Al matrix (Fig. 2a). EDS analysis at the exhibited points A and B reveals that the small dark areas in the matrix are B-rich, while the white fragments are Ti-rich remnants. According to the thermodynamic data, TiB2 is the most stable phase in Al-Ti-B system. Nevertheless, it seems that the presence of a large fraction of Al can limit the inter-diffusion of Ti-B in the powder mixture; preventing the completion of TiB2 synthesis. Suppressive effect of large Al amount on Ti-B reaction in Al-Ti-B system has been reported previously [24]. Typical cross-sectional SEM micrograph of TiB2 20 wt%Al particles after a total milling time of 20 h shows the formation of fine particles which are, according to EDS point analysis at the depicted point C, TiB2 particles (Fig. 2b). It can be concluded that the diffusion path between Ti and B decreases with decreasing the Al content, which can assist the reaction between them and formation of TiB2. SEM image of MAed TiB2 10 wt%Al particles (Fig. 2c) indicates that it consists of TiB2 particles with a lower amount of Al as the binding element. Moreover, the size of TiB2 particles seems to be larger in this powder compared to that in TiB2 20 wt%Al composite powder. Since the Al binder can facilitate the distribution of TiB2 particles during ARB, the TiB2 20 wt%Al composite powder obtained from 20 h of MA seems to be the appropriate powder for ARB. The Ti-B reaction is completed in this powder mixture with the highest possible Al
content. Fig. 3 shows the morphology of powders utilized for ARB process. The size range of as-received TiB2 powder is roughly less than 10 μm with mostly angular particles (Fig. 3a). After milling with 20 wt% Al powder for 5 h, a mixture of fine angular particles is achieved with a size of less than 2 μm (Fig. 3b). The in-situ TiB2 20 wt%Al powder obtained by 20 h of MA shows nearly spherical submicron particles with many particles having the size less than 500 nm (Fig. 3c). 3.2. ARB processing 3.2.1. Microstructural evolutions To investigate the bonding development between layers and the distribution of particles during ARB, the microstructure of AlTiB2 (as-received) composite after various ARB cycles is selected as typical and the results are presented in Fig. 4. After the first and the third cycle (Fig. 4a and b), the particles are seen at layer interfaces and a non-uniform distribution of large clusters of TiB2 particles as well as particle free zones can be observed. After five ARB cycles (Fig. 4c), particles are removed from interfaces to some extent, but the distribution of TiB2 particles in Al matrix seems still non-uniform. As the ARB is continued up to seven cycles (Fig. 4d), the interface between Al layers cannot be easily detected and a nearly uniform distribution of TiB2 particles in Al matrix is obvious. In this stage, the particles are separated from layer interfaces and moved into the matrix. This can be explained by film
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Fig. 4. Low magnification SEM micrographs of the Al-TiB2 (as-received) composite after various ARB cycles; (a) one, (b) three, (c) five and (d) seven cycles.
theory as follows [30]. During ARB process, two opposing work hardened layers or brittle surface oxide layers on Al strips break up coherently and the underlying metals extrude through the produced channels between the cracks from both sides of the interfaces. The presence of particles at interfaces can be considered as a preventing barrier for the film theory [31]. However, the Al matrix deforms plastically through the particle clusters and fresh metal surfaces can be obtained during ARB. At the same time, the rolling pressure can generate sufficient force and good bonding at layer interfaces even in the presence of ceramic particles. At the same time, with increasing the ARB cycles, the number of Al layers increases while their thickness decreases, leading to the further distribution of the particles and their movement from the interfaces into the bulk of the Al matrix. It is noted that the low magnification SEM image of the composite (as presented in Fig. 4) is not appropriate to show the bonding quality of the individual particles with the Al matrix. It just explains the fragmentation of the particles at interfaces and their distribution in the Al matrix. High magnification SEM micrographs of Al-TiB2 (in-situ) composite are presented in Fig. 5 to indicate the structure of TiB2 fragments and the quality of bonding between the particle and matrix after different ARB cycles. After the first ARB cycle, the material is severely porous at the interfaces with a weak bonding between the particle and matrix. These pores may be between either in the layer interface or near the Al and the particle interfaces. With increasing the number of cycles, the amount and size of porosities decreases and bonding at interfaces improves significantly. Fig. 6 presents the SEM micrographs of various Al-TiB2 composites on RD-TD plane after three and seven ARB cycles. In all composites, clusters of TiB2 particles as well as particle free zones can be seen apparently after three ARB cycles. The agglomeration of particles is reduced to some extent after seven ARB cycles. It has
been stated that the distance between the particles inside the clusters can increase due to the elongation of the clusters in rolling direction [8,32]. Consequently, the dense clusters are changed into diffuse ones with islands of Al matrix between them [33]. The microstructure of Al-TiB2 composites in the as-received and premixed conditions reveals that by premixing of TiB2 particles with Al, improved distribution with less agglomeration of TiB2 particles can be achieved after each ARB cycle. It can be concluded that the existence of Al binder between TiB2 particles in the mixed Al-TiB2 powder can facilitate the extrusion of Al matrix through initial clusters of particles. The composite with the in situ Al-TiB2 powder exhibits the best distribution as well as the least agglomeration among the others after the third and the seventh cycle. It has been reported that the in-situ synthesis results in clean and defect free interfaces between the in-situ reinforcing particles and the adjacent phase. This can provide a strong interfacial bonding between the particles and the matrix and can improve the mechanical properties of the composite [22]. TEM micrographs and the corresponding selected area diffraction (SAD) patterns (take with 1.3 mm aperture size) on RD-TD plane of Al-TiB2 (in situ) composite are shown in Fig. 7. After three ARB cycles, the grains are divided into subgrains through formation of the dislocation walls. SAD pattern of the sample shows a near net pattern, indicating that the boundaries have low misorientation in the corresponding area. Contours or fringes that can be seen mostly near the boundaries reflect the internal stress existing in grains. After seven ARB cycles (Fig. 7b), a significant decrease in grain size is observed and the microstructure consists of an array of ultrafine equiaxed grains with a less regular arrangement. SAD pattern shows the ring like patterns consisting of separate spots that indicates there are higher fractions of boundaries with high angles of misorientation. However, the size of grains seems to vary in different areas that can be attributed to the
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Fig. 5. SEM micrographs of Al-TiB2 (in situ) composite after various ARB cycles; (a) three, (b) five, (c) seven cycles and d) EDS analysis of point A.
distribution of TiB2 particles within the matrix. Generally, the formation of the ultra-fine structure in the highly strained materials is regionally inhomogeneous [34] because of the friction between the rolls and the strip surfaces. The sequence of grain refinement during SPD process has been explored by several researchers as follows [35]. A fine dislocation cell structures with low-angle grain boundaries is created by arrangement of dislocations into the low energy dislocation structure [36]. The dislocation cell structure or subgrains become finer and their misorientation angle increases due to operation of multi-directional slip and absorption of further dislocations [37]. Finally, a distinct grain structure with distinguishable high-angle boundaries is generated when the sample is subjected to a very intense plastic strain. Fig. 8 shows the STEM micrographs of Al-TiB2 (in situ) composite after three and seven ARB cycles. After three cycles of ARB, the structure contains grains with average size of about 850 nm and includes the subgrains with dislocation cell structures. The dislocations can be observed in grains with different geometries. Some of them accumulate as tangles of multiple dislocations and some emerge as individual dislocations in coarse grains (indicated in Fig. 8a). The contrast change in some grains can be considered as an evidence for the initiation of subgrain formation. The presence of subgrains within the grains is an indication for mechanism of continuous dynamic recrystallization mechanism [38]. This mechanism is based on continuous absorption of dislocations in subgrain boundaries, followed by the formation of new grains separated by high angle boundaries. After seven ARB cycles, the average grain size is decreased to about 300 nm. The dislocation density inside the grains is very low in comparison with that
observed after three ARB cycles. It has been stated that the development of fine structures during SPD is accompanied by a decrease in dislocation density at large strains. Indeed, severely deformed materials commonly contain fine grains that are completely free of dislocations [35]. Extinction of dislocations inside the grains in the final stages of ARB has been reported by previous researcher [39]. 3.2.2. Mechanical properties Fig. 9a illustrates the representative engineering stress–strain curves of the annealed Al and Al-TiB2 (in situ) composite after different cycles of ARB processing. In general, the tensile test is not an appropriate tool to analyze the elastic properties of the materials and more accurate test methods are required. Therefore, the engineering stress-strain curves obtained from tensile test do not exhibit a significant difference between the elastic modulus of the composite materials produced by ARB. A similar trend can also be observed in literature such as in Al/SiC [7], Al/CuP [40] and Al/ Al2O3/SiC composite [41]. The variation in the ultimate tensile strength (UTS) and uniform elongation with the number of cycles is also plotted in Fig. 9b. After the first cycle, a significant increase in UTS to about 155 MPa is observed in comparison to that of the annealed Al (47 MPa). The UTS increases gradually with increasing the number of ARB cycles up to 171 MPa after seven ARB cycles. Tensile curves of the ARB processed composite show immediate UTS after yielding unlike the annealed sample. Contrary to the strength, the elongation decreases drastically after the first ARB cycle. During tension test, the response of nano/ultra-fine grained polycrystalline metals is different from that of coarse grain metals.
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Fig. 6. SEM micrographs of various composites; a), c), e) three cycles and b), d), f) seven cycles.
Several reports show that the work hardening is absent in ultrafine grained materials that are produced by severe plastic deformation [42]. This leads virtually to necking at the yield stress (in tension), and the net results is a low tensile elongation. In the other words, the yield stress and UTS of SPD processed materials are almost identical. Therefore, the variation of UTS with the number of ARB cycle (that is presented in Fig. 9b.) can be considered as a good approach to show the effect of ARB on the yield stress. Two important mechanisms that are responsible for the deformation behavior of the SPD processed materials are strain hardening and grain refinement. The former is dominated in the
early stages and the latter is occurred in the last stages of SPD [43]. The increase in dislocation density within the grains and formation of low angle dislocation boundaries are the main reasons of strain hardening in the early stages of ARB. However, in the final stages, the size of dislocation boundaries decreases, the misorientation at these boundaries increases progressively and grain refinement to the ultra-fine scale causes the increase in the strength [44]. In addition, the existence of TiB2 reinforcing particles can also be responsible as additional barriers for moving dislocations and consequently increase in the tensile strength. Considering the Orowan strengthening mechanism that involves dislocations bowing around the particles, dispersion of fine
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(a)
(b)
Fig. 7. TEM micrographs and corresponding SAD patterns of Al-TiB2(in situ) composite after a) three and b) seven ARB cycles.
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Fig. 8. STEM micrographs of Al-TiB2(in situ) composite after a) three and b) seven ARB cycles.
180
180
90 60
Annealed Al
30 0 0.00
0.10
0.20 0.30 0.40 Engineering Strain (mm/mm)
Ultimate Tensil Stress (MPa)
200
150
30
Al-TiB2(In Situ) Al-TiB2(Mixed) Al-TiB2(As-received)
120
90
60 0.0
0.50
0.1
0.2 0.3 0.4 Engineering Strain (mm/mm)
0.5
Fig. 10. The engineering stress–strain curves of final Al-TiB2 composites obtained from seven ARB cycles.
25 150 20 100
15
Ultimate strength Uniform Elongation
10
50
Uniform Elongation (%)
Emgineering Stress (MPa)
120
Engineering Stress (MPa)
cycle7 cycle 5 cycle3 cycle1
150
5 0
0
1
2
3
4
5
6
7
0
Number of ARB cycle Fig. 9. a) The engineering stress–strain curves of annealed Al and Al-TiB2 (in situ) composite after different ARB cycles and b) variation of tensile strength and elongation with respect to the number of ARB cycles.
particles is an effective factor for the strengthening. The extreme decrease in the elongation can be explained by the increase in density of dislocation at the first ARB stages that promotes the nucleation of cracks. Moreover, TiB2 reinforcing particles can induce voids and cracks that can reduce the ductility of the composite. However, a slight increase in elongation after successive ARB cycles can be seen which can be resulted from enhancement
of the bonding at interfaces. In order to compare the deformation behavior of Al-TiB2 composites, the engineering stress–strain curves of all composites (after seven ARB cycles) are presented in Fig. 10. The Al-TiB2 (in situ) composite shows the largest UTS (171 MPa) while the least UTS belongs to Al-TiB2 (as-received) (156 MPa). Different characteristics of the reinforcing particles such as morphology, size, quantity and distribution can affect the strength of the composites. Since the Al-TiB2 (in situ) composite reveals a more uniform distribution of fine TiB2 particles with less porosity and spherical morphology, the highest UTS is expected for this composite. Defects such as porosity at particle-matrix interface and agglomeration of particles are considered as main reasons for composite failure. Fig. 11 presents typical fracture surfaces of Al-TiB2 (in situ) composite after three and seven ARB cycles. The fracture surfaces show a typical ductile fracture that is characterized by dimples. This mode of fracture occurs by the formation and coalescence of microvoids, crack propagation and final fracture [45]. Several voids can be elongated in one direction because of unequal triaxial stresses [46]. The presence of particles at the bottom of some dimples, indicated in the micrographs, can be considered as an evidence for the nucleation of voids from particle-matrix interfaces [47]. The agglomeration of particles after the third cycle
M. Askarpour et al. / Materials Science & Engineering A 677 (2016) 400–410
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Fig. 11. SEM micrographs of the fracture surfaces after tensile experiment of Al-TiB2 (in situ) composite obtained from (a) three and (b) seven ARB cycles.
(Fig. 11a) is obvious on the fracture surface while the no particle clustering can be seen after the seventh cycle (Fig. 11b). The presence of the particles can change the morphology of dimples in a manner that the dimples nucleated at the particle sites (as indicated in Fig. 11b) are deeper and larger than that nucleated in other regions. These results are consistent with those obtained in the previous sections (Section 3.2.2) where a uniform distribution of particles was observed after seven ARB cycles. It is noted that the dimples are no longer apparent at regions where the particles clustering occur; indicating that agglomeration of particles can affect the mode of fracture as well as mechanical properties.
5. The fracture surface of Al-TiB2 (in situ) composite shows several dimples, which is an indication of ductile-type fracture. The agglomeration of particles on the fracture surface of the composite after the third cycle is more evident.
Acknowledgment Financial support provided by Shahid Chamran University of Ahvaz (Grant no. 94-3-02-31579) is gratefully acknowledged. Authors would like to thank Mr. E. Bagherpour for the preparation of TEM samples and taking TEM images in Metallic Materials Science Laboratory of Doshisha University, Japan.
4. Conclusions Al-TiB2 metal matrix composites were fabricated by ARB up to seven cycles. The reinforcing particles were prepared in three different processing conditions: as-received TiB2, mixed TiB2-Al and in-situ synthesized TiB2-Al. The most important conclusions are: 1. In-situ synthesized TiB2-Al powder with 10, 20 and 30 wt% Al is prepared by 20 h MA. It is concluded that the in-situ synthesized TiB2 20 wt%Al can be the appropriate powder for ARB because of completing the Ti-B solid state reaction. The higher inter-diffusion between Ti and B in the powder mixture (compared with TiB2 30 wt%Al) and the smaller size of TiB2 particles (compared with TiB2 10 wt%Al) are the advantages of TiB2 20 wt%Al powder mixture. 2. In comparison to the Al-TiB2 (as-received) and Al-TiB2 (mixed) composites, the best distribution as well as the least agglomeration is obtained in the Al-TiB2 (in-situ) composite. It was attributed to the formation of the clean, defect free interfaces between the in-situ produced powders and the adjacent phase. 3. After three ARB cycles, a subgrain structure with an average grain size of about 850 nm is obtained in the Al-TiB2 (in-situ) composite. ARB processing up to seven cycles leads to an ultrafine grained structure with an average size of about 300 nm in the Al-TiB2 (in-situ) composite. 4. The highest tensile strength of about 171 MPa is achieved for the Al-TiB2 (in situ) composite because of a more uniform distribution of fine TiB2 particles with less porosity and spherical morphology.
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