reinforcement interfaces and properties of Al-TiB2 composites

Accepted Manuscript Effect of cryomilling on matrix/reinforcement interfaces and properties of Al-TiB2 composites Özge Balcı, Duygu Ağaoğulları, Hasan Gökçe, M.Lütfi Öveçoğlu, Mehmet Somer PII:

S0925-8388(18)31783-3

DOI:

10.1016/j.jallcom.2018.05.098

Reference:

JALCOM 46071

To appear in:

Journal of Alloys and Compounds

Received Date: 1 June 2017 Revised Date:

7 May 2018

Accepted Date: 8 May 2018

Please cite this article as: Ö. Balcı, D. Ağaoğulları, H. Gökçe, M.Lü. Öveçoğlu, M. Somer, Effect of cryomilling on matrix/reinforcement interfaces and properties of Al-TiB2 composites, Journal of Alloys and Compounds (2018), doi: 10.1016/j.jallcom.2018.05.098. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

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ACCEPTED MANUSCRIPT Effect of cryomilling on matrix/reinforcement interfaces and properties of Al-TiB2 composites Özge Balcı1,2*, Duygu Ağaoğulları3, Hasan Gökçe4, M. Lütfi Öveçoğlu3, Mehmet Somer1,2 1

Koç University, Department of Chemistry, Rumelifeneri Yolu, 34450 Sarıyer, Istanbul,

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Turkey Koç University Akkim Boron-Based Materials and High Technology Chemicals Research

and Application Center, Rumelifeneri Yolu, 34450 Sarıyer, Istanbul, Turkey 3

Istanbul Technical University, Faculty of Chemical and Metallurgical Engineering,

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Department of Metallurgical and Materials Engineering, Particulate Materials Laboratories (PML), Ayazağa Campus, 34469 Maslak, Istanbul, Turkey 4

Istanbul Technical University, Prof. Dr. Adnan Tekin Materials Science and Production

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Technologies Applied Research Center, Ayazağa Campus, 34469 Maslak, Istanbul, Turkey *Corresponding Author: [email protected], Tel.: +90-212-338-0933; Fax: +90-212-338-1548

Abstract

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The present study reports the effects of cryomilling on the microstructural and mechanical properties of pressurelessly sintered Al-xTiB2 (x = 5, 10, or 15 wt.%) composites. Composite powders milled under cryogenic conditions in the presence of liquid nitrogen were compacted

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in a hydraulic press, and the compacts were pressurelessly sintered at 650 ºC for 5 h under Ar/H2 gas. The effects of cryomilling time (10 or 20 min) and TiB2 content (5, 10, or 15

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wt.%) on the microstructures, densities, and mechanical properties (wear characteristics, hardness, elastic modulus, indentation response) of the sintered products were investigated. Because the milling process was under cryogenic conditions, the formation of intermetallic phases between the matrix and reinforcement particles, such as Al3Ti, was not observed at the interfaces of the sintered Al-TiB2 composites, even when 15 wt.% TiB2 was used. Cryomilling for 10 min followed by pressureless sintering of Al matrix composites reinforced with 5–15 wt.% TiB2 particles resulted in bulk samples with relative densities of 97.38–

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ACCEPTED MANUSCRIPT 98.43%, hardness values of 0.75–1.35 GPa, wear volume losses of 0.276–1.265 mm3, and elastic moduli of 283–433.7 GPa. Keywords: Metal matrix composites; Powder metallurgy; Microstructure; Mechanical

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properties; X-ray diffraction; Scanning electron microscopy 1. Introduction

Al-based metal matrix composites (MMCs) reinforced with continuous/discontinuous

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reinforcements have attracted considerable attention in various areas, such as aerospace and automotive industries, particularly because of their high specific moduli and strengths, as well

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as high wear and creep resistances [1–4]. Several technologically attractive materials can be fabricated by combining ductile Al-based matrices with different types of high-strength reinforcements to produce lightweight composites with improved specific strengths [3–4]. In simple terms, by creating suitable combinations of a matrix, reinforcement material, and

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fabrication process, the increasing demand for different industrial applications can be met [4– 5]. Ex situ or in situ reinforced Al-based MMCs have been manufactured using different techniques such as reactive squeeze casting and rapid solidification processing, exothermic

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9].

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dispersion, reactive hot pressing, combustion synthesis, and powder metallurgical routes [4–

Typical reinforcement compounds used in Al-based MMCs are oxides (e.g., SiO2 or Al2O3) [8], carbides (e.g., SiC or TiC) [9–11], nitrides (AlN) [12], and borides (e.g., TiB2 or ZrB2) [13–14], as well as metallic glasses [15] and complex metallic alloys in the form of fibers, flakes, or particulates [11, 15]. Among them, particulates are very attractive because particulate-reinforced Al matrix composites (AMCs) are relatively easy to produce, and the costs of the raw materials are lower than those of Al alloy matrix (e.g., Al-Si or Al-Cu) composites reinforced with fibers or flakes [4, 8, 11, 14–16]. In addition, there is increasing

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ACCEPTED MANUSCRIPT interest in the use of TiB2 particles as the reinforcement material owing to their attractive properties such as their high melting point, high hardness, superior wear resistance, high electrical conductivity, and considerable chemical and thermodynamic stability in Al matrices [17–21]. In particular, the incorporation of TiB2 ceramics into Al matrices offers the

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advantage of enhanced tensile strength, fracture toughness, and wear resistance compared to monolithic Al [14, 22–26].

AMCs reinforced with TiB2 particles have previously been fabricated via direct melt reaction,

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hot isostatic pressing, or friction stir welding processes [27–30]. Although these methods

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suffer from some considerable problems such as agglomeration and inhomogeneous distribution of TiB2 particles in the matrices [18, 31], the microstructural and mechanical properties of Al-TiB2 composites were successfully improved by means of several techniques recently reported in the literature [32–38]. Powder metallurgy through solid-state sintering, on the other hand, is particularly suitable for the fabrication of composites, because it enables a

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great degree of microstructural control, for example, the size, morphology, and volume fraction/distribution of the reinforcements [11, 39]. For this reason, the role of powder preparation on the properties of Al-TiB2 composites has been investigated in a few studies

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[40–41]. Nonetheless, some undesirable interfacial reactions have been reported between the

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matrix and the reinforcing particles. For example, during the fabrication of Al-TiB2 composites, several researchers have reported the formation of intermetallic Al3Ti, which considerably reduced some mechanical properties, including wear resistance and fatigue life [14, 18, 31, 42]. Therefore, preventing the formation of this undesirable Al3Ti phase is of interest when the powder metallurgical route is used. Mechanical alloying (MA) as a composite powder preparation technique has been employed to minimize undesirable reactions between the matrix and reinforcement, offering the advantage of room temperature manufacturing [3, 43]. However, in Al-TiB2 composites obtained via MA followed by 3

ACCEPTED MANUSCRIPT sintering processes, small amounts of Al3Ti flakes have also been observed [14]. In addition to MA, cryomilling (CM) is a powder preparation technique in which high-energy milling is performed in cold environments such as liquid nitrogen [44]. Owing to the extremely low temperatures in the cryogenic media, CM enables a reduction of milling time, grain size, and

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particle agglomeration [44–45]. In addition, it can eliminate undesirable reactions during the milling and transfer of high energy to the powders [46–48]. From this point of view, CM before sintering seems to be a promising powder preparation method for AMCs that are

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reinforced with TiB2 particles. However, such studies have not yet been reported.

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To overcome the problem of interfacial reactions between the matrix and the reinforcement, a new route is proposed in the present study for producing TiB2 particulate-reinforced AMCs via CM followed by pressureless sintering. The role of interfaces with respect to CM on the microstructures of the composites was evaluated. Furthermore, the effects of the TiB2 content

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and CM time on the microstructure and mechanical properties were also investigated. 2. Experimental procedures 2.1. Powder preparation

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Elemental Al powders (Alfa AesarTM, 99.5% purity, 12 µm average particle size) as the matrix and TiB2 powders (Alfa AesarTM, 99.5% purity, 40 µm average particle size) as the

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reinforcement were used to prepare the composite powders. Al and TiB2 powders were mixed in a WABTM T2C Turbula blender for 30 min to derive compositions of Al-5 wt.% TiB2, Al10 wt.% TiB2, and Al-15 wt.% TiB2 (hereafter termed as Al5TiB2, Al10TiB2, and Al15TiB2, respectively). The blends were cryomilled (CM’d) for 10 or 20 min to prepare the composite powders. The rationale for using CM times of 10 and 20 min is based on the unavoidable Fe contamination in the powders that originates from the milling media; this contamination is considerable for milling times longer than 20 min. CM was conducted in a Spex™ 6870 Freezer/Mill (at a rate of 600 collisions/min) using a cylindrical polycarbonate vial and 4

ACCEPTED MANUSCRIPT stainless steel rods. Cryogenic conditions were provided using liquid nitrogen (Linde™, refrigerated) externally circulated around the milling vial. Loading, sealing, and unloading of the milling vials were conducted under Ar gas (Linde™, 99.999% purity) in a Plaslabs™ glove box.

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2.2. Sintering

After CM, each powder mixture was compacted in a 10-ton MSETM MP-0710 uni-action hydraulic press under a uniaxial pressure of 300 MPa to obtain a cylindrical preform with a

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diameter of 12.7 mm. For tensile strength measurements, each powder mixture was compacted to obtain a regular tensile test specimen (overall length: 90 mm, distance between

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shoulders: 48 mm, grip section: 21 mm, width of grip section: 9 mm, length of reduced section: 25 mm, gauge width: 5.5 mm, gauge length: 20 mm) using the same conditions as those of the cylindrical preform. All compacts were sintered at 650 ºC for 5 h in a Linn™ HT1800 high-temperature controlled atmosphere furnace at a heating and cooling rate of 10

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ºC/min. To prevent probable surface oxidation of the compacts, the sintering process was conducted under Ar/H2 gas flow conditions. Sintered samples fabricated from the composite powders prepared by CM for 10 or 20 min are hereafter termed as AlxTiB2-10 and AlxTiB2-

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20 (x = 5, 10, 15), respectively. 2.3. Characterization

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X-ray diffraction (XRD) investigations of the composite powders and sintered samples were carried out using a BrukerTM D8 Advanced Series powder diffractometer with CuKα (λ = 1.5406 Å) radiation in the 10–90º 2θ range incremented with a step size of 0.02º at a rate of 2º/min. International Center for Diffraction Data (ICDD) powder diffraction files were utilized for the identification of crystalline phases. The average crystallite sizes of the Al phases in the CM’d composite powders were predicted using BrukerTM-AXS TOPAS V3.0

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ACCEPTED MANUSCRIPT software. Fe contaminations in representative samples were determined by X-ray fluorescence (XRF) using a Bruker Tiger S8 XRF spectrometer. For characterization of the bulk samples, cylindrical specimens with diameters of 12.7 mm and lengths of 5 mm were used. To obtain a scratch-free mirror finish for microstructural

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analyses, indentation measurements, and wear tests, at least two specimens for each sintered sample type were subjected to a typical metallographic preparation procedure. Specimens were hot-mounted in a StruersTM LaboPress-1 and ground and polished in a StruersTM

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Tegrapol-15 instrument. Backscattered electron images of the bulk samples were obtained using a HitachiTM TM-1000 scanning electron microscope (SEM) operating at 15 kV.

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Secondary electron SEM images and energy-dispersive X-ray spectroscopy (EDX) analyses of selected samples were achieved using a ZeissTM Ultra Plus field emission SEM equipped with an EDX detector. Sample densities were measured in ethanol using the Archimedes method, and the results were reported as the arithmetic mean of three replicate measurements

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for each of the two specimens for each sample type.

All sintered samples were subjected to reciprocating sliding wear tests at room temperature in a laboratory atmosphere by a Tribotech™ Oscillating Tribotester using a 6-mm alumina ball

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under an applied force of 3 N, with a sliding speed of 5 mm/s and a stroke length of 2 mm for a total sliding distance of 25 m. Worn surfaces screened with a LeicaTM ICC50 HD optical

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microscope (OM) were examined using a Veeco™ Dektak 6 M Stylus profilometer: Changes in penetration depth and horizontal distance of the sintered samples were reported. The wear test results, in terms of wear volume loss values, were the arithmetic mean of three replicate measurements for each sample. Relative wear resistance values for all samples were calculated as the ratio of the highest measured wear volume loss to the wear volume loss of the individual samples. Wear tracks of selected samples were imaged using a ZeissTM Ultra Plus field emission SEM.

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ACCEPTED MANUSCRIPT Depth-sensing indentation measurements of the samples were taken at a peak load of 50 mN using an Agilent G200 nanoindenter. Load–displacement curves of the sintered samples were obtained from the Al matrices and TiB2 reinforcements without a holding time at the peak load. In addition, the hardness values and elastic moduli of the matrices and reinforcements,

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obtained from nanoindentation tests of the sintered samples, are reported. For comparison of the calculated hardness values of the composites with the experimental data, Vickers microhardness values were determined for representative samples using a Shimadzu™ HMV

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microhardness tester under a load of 25 g for 10 s. The microhardness test result for each sample represents the arithmetic mean of ten successive indentations and the standard

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deviation. Tensile strength measurements were taken on tensile specimens using a ShimadzuTM Autograph AGS-J table-top type universal tester with capacity of 10 kN to obtain their stress–strain curves. Each point at the stress–strain curve of each sample included the arithmetic mean of three different measurements.

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3. Results and discussion

Figs. 1(a)–(c) and (d)–(f) illustrate the XRD patterns for the Al5TiB2, Al10TiB2, and Al15TiB2 powders CM’d for 10 min and 20 min, respectively. All CM’d powders contained

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Al (ICDD Card No: 04-0787; Bravais lattice: face-centered cubic, a = 4.0494 Å) and TiB2 (ICDD Card No: 35-0741; Bravais lattice: primitive hexagonal, a = 3.0303 Å, c = 3.2295 Å)

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phases. A slight increase in the intensities of the TiB2 peaks can be observed for both milling times (Figs. 1(a)–(c) and (d)–(f)) as the TiB2 content changed from 5 to 15 wt.%. No significant broadening due to the extended CM time (from 10 to 20 min) occurred in the XRD peaks of the Al and TiB2 phases. Cold welding, fracturing, and rewelding mechanisms during milling processes can cause degradation of the particles and trigger the formation of intermetallic phases [43, 49]. However, CM carried out in the presence of externally circulated liquid nitrogen did not induce the formation of secondary phases. In addition, no

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ACCEPTED MANUSCRIPT impurities, worn or torn from the milling media, were observed within the detection limit of the utilized diffractometer (> 2 wt.% of the sample). Fe contamination originating from the milling media, on the other hand, did occur, at levels of 4.4 and 6.2 ppm, respectively, for the Al10TiB2 powders CM’d for 10 and 20 min, determined by XRF analyses. Fe impurities were

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found to be negligible after a subsequent sintering processes for the powders CM’d for 10 and 20 min. CM for 30 min resulted in an increase in the amount of Fe impurities to 11.5 ppm for the Al10TiB2 powders, justifying the choice for the utilized lower CM times.

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A benefit of the CM process is in grain refinement at a much higher rate with a shorter milling time than with conventional milling [50]. Based on this, the average crystallite sizes of the Al

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phases in the 10- and 20-min CM’d Al5TiB2, Al10TiB2, and Al15TiB2 powders were measured, as shown in Table 1. There was a remarkable decrease (> 35%) in the average crystallite size of the Al phase as the CM time increased from 10 to 20 min. It was earlier reported that mechanically alloyed AlxTiB2 (x = 5, 10, and 15 wt.%) powders had an average

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crystallite size of approximately 175 nm after 1 h of milling [14]. Even when 10 min of CM was applied under cryogenic conditions, the powders here had lower average crystallite sizes (≤ 168 nm) than the mechanically alloyed powders. Furthermore, the average crystallite size

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decreased as the TiB2 content increased from 5 to 15 wt.% owing to the hard and brittle character of TiB2 and the cryogenic milling environment.

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Figs. 2(a)–(f) show the XRD patterns of bulk samples sintered from the Al5TiB2, Al10TiB2, and Al15TiB2 powders CM’d for 10 and 20 min. As seen in Figs. 2(a)–(c), only Al and TiB2 are present after the sintering. The absence of secondary phases after CM for 10 min and subsequent sintering conforms well with the theory of using chemically stable ceramic reinforcement particles in light-metal matrix systems [1, 42, 51–52]. According to the XRD patterns in Figs. 2(d)–(f), however, the extended CM time of 20 min resulted in an Al2O3 phase (ICDD Card No: 46-1212; Bravais lattice: primitive rhombohedral, a = 4.7587 Å, c =

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ACCEPTED MANUSCRIPT 12.9929 Å) for the Al5TiB2, Al10TiB2, and Al15TiB2 samples after sintering. However, the XRD peak of the Al2O3 phase is more developed for the Al15TiB2 sample (Fig. 2(f)) than for the Al5TiB2 (Fig. 2(d)) and Al10TiB2 (Fig. 2(e)) samples. Although the emergence of this Al2O3 phase was not expected considering the initial materials and CM and sintering

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conditions, a very small amount of Al2O3 phase in the 20-min CM’d and sintered samples may be related to oxygen adsorption by activated Al particle surfaces after prolonged milling and to the reaction of Al and O during sintering. It can therefore be stated that an increase in

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CM time is not beneficial to obtain contamination-free reinforcement and matrix interfaces, because it creates more reactive surfaces that are open to any possible reaction during

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handling of the powders. EDX measurements of the Al5TiB2, Al10TiB2, and Al15TiB2 samples CM’d for 20 min and later sintered showed 0.38, 0.65, and 0.98 wt.% O, respectively. These oxygen contents correspond to approximately 0.8, 1.4, and 2.1 wt.% Al2O3 phase in the Al5TiB2, Al10TiB2, and Al15TiB2 samples, respectively. Although Al2O3

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was only formed after the sintering, it could contribute to the hardness and mechanical properties as a reinforcing agent. In some studies related to the production of AMCs CM’d in a liquid N2 environment, some secondary phases have also been observed after the sintering,

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such as AlN and Al(ON) [48]. However, CM carried out with liquid N2 externally circulating around the milling vial in this study inhibited the formation of such nitride- and oxynitride-

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based compounds. In addition, no Fe impurities originating from the milling vial/balls were detected after the sintering, which could mean that the amount was < 2 wt. %. The formation of brittle reaction products—such as the Al3Ti phase—among the reinforcement and matrix interfaces has been encountered in some investigations in which casting and powder metallurgy processes were utilized for the fabrication of Al-TiB2 composites [18, 53–54]. It was previously reported that the formation of Al3Ti after mechanical alloying and sintering is due to partial bond weakening of TiB2 by the high increase in impact energy and the

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ACCEPTED MANUSCRIPT temperature caused by the continuous collisions and partial melting of Al at the sintering temperature [14]. The lack of the intermetallic phase Al3Ti in the 10- and 20-min CM’d and sintered samples may be attributed to the special feature of the applied milling process in which externally circulating liquid nitrogen did not allow the temperature increase and

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absorbed the released excessive energy. Thus, the limitations in diffusion processes at cryogenic temperatures provided an advantage for CM over the room temperature milling processes [50].

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Figs. 3(a)–(f) represent the backscattered electron SEM images of bulk samples sintered from the Al5TiB2, Al10TiB2, and Al15TiB2 powders that were CM’d for 10 and 20 min. For both

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CM times, the increase in the amount of TiB2 and its homogeneous distribution throughout the microstructure can be clearly seen. The main problems encountered in casting processes, such as a lack of well-distributed reinforcement particles, the formation of nonreinforced or clustered areas, and the occurrence of segregation, were not observed in the microstructures of

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the CM’d and sintered samples. According to the earlier performed studies, Al3Ti phases in the form of flakes occur in TiB2 particulate-reinforced Al matrices [14, 18, 54]. Here, in contrast, no flake formation was detected in the SEM images in Figs. 3(a)–(f). This absence of

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Al3Ti flakes was further supported by the XRD study (Fig. 2) of the samples. Additionally, the sintered specimens that were CM’d for 20 min had higher porosities than those CM’d for

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10 min, which may be related to the presence of small amounts of Al2O3 phase. CM enabled the fabrication of sintered samples without the formation of secondary phases from the Al and TiB2 reactants. The Al5TiB2, Al10TiB2, and Al15TiB2 composites produced in the present study are thus free of some of the drawbacks, such as the formation of thick Al3Ti flakes, fragmentation of thin Al3Ti flakes, occurrence of cracks and voids in the external surfaces of the Al3Ti flakes, and decrease of densification rate. To study the clean interface between the Al matrix and TiB2 reinforcement particles, a secondary electron SEM

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ACCEPTED MANUSCRIPT image and EDX maps of the matrix/reinforcement interface were taken from Al10TiB2-10 as a representative sample (Fig. 4). The Ti and B elemental maps overlap almost entirely, suggesting the presence of TiB2 particles, whereas the Al elemental map corresponds to the whole area surrounding the TiB2 particles. There is no indication of a secondary phase or

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contaminations at the interface of the matrix and particles. Furthermore, the secondary electron SEM image and EDX point analyses (points A, B, and C) of the sintered Al15TiB220 sample matrix/reinforcement interface (Fig. 5) show light gray particles whose EDX

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analysis at point A indicates pure Al representing the matrix. B shows white, irregularly shaped particles accumulated in the interface region. EDX analysis taken from point B

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indicates the presence of Al and O that stem from the compound Al2O3, which has also been clearly identified in the XRD pattern (Fig. 2(f)). The EDX measurement on the particulate reinforcement labelled as C in the SEM image showed Ti and B as majority components together with Al (0.56 at.%) from the matrix. The occurrence of Al2O3 contamination in the

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20 min of CM’d and sintered samples, as well as in the deposited region in the composite, were also proved by the microstructural analyses. The relative density–TiB2 content curves for the bulk samples sintered from Al5TiB2,

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Al10TiB2, and Al15TiB2 powders CM’d for 10 and 20 min are shown in Fig. 6. For comparison purposes, the relative densities for samples sintered from mechanically alloyed

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(MA’d) powders are also included in Fig. 6 (adapted from [14]). As seen in Figs. 6(c) and (d), Al10TiB2 samples had higher relative density values (98.43 and 96.76%) than Al5TiB2 and Al15TiB2 samples, for both CM’d for 10 or 20 min. Similarly, among AlxTiB2 (x = 5, 10, and 15 wt.%) samples MA’d for 8 h and sintered, Al10TiB2 (Fig. 6(a)) had the highest density value of 94.66% [14]. This may be an indication that, by increasing the TiB2 content, densification occurs up to a limited value in the Al matrix. It should also be noted that, even at the same TiB2 content (10 wt.%), a change in the milling type before sintering from MA to

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ACCEPTED MANUSCRIPT CM resulted in a higher densification, which depended on the efficiency of particle size reduction without releasing contamination (Figs. 6(a)–(d)). Overall, as clearly seen from Figs. 6(a)–(d), CM provided higher densities than MA at all TiB2 contents. According to Figs. 6(c) and (d), all samples CM’d for 10 min had higher relative density values than those CM’d for

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20 min. Although 20 min of CM time enabled a higher reduction in crystallite size for the powders, increased densification could not be achieved after sintering. This result may be related to the presence of pores in the sintered bodies (Figs. 3(d)–(f)), which were caused by

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the formation of the Al2O3 phase after sintering the 20-min CM’d powders (Figs. 2(d)-(f) and Fig. 5). In this case, the sintering process was not sufficient to achieve bonding of both the

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TiB2 reinforcement and unavoidable byproduct Al2O3 with the Al matrix and overcome the large number of defects created by CM [48]. Based on this, the unexpected Al2O3 formation decreased the densification rate of the composites.

The changes in penetration depth and horizontal distance (obtained from the wear profiles) for

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the sintered samples are shown in Fig. S1 in the supplementary data. Among all samples, the deepest penetration and longest horizontal distance were observed in the Al5TiB2-10 sample. At the same CM time of 10 min, both the penetration depth and horizontal distance

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dramatically decreased as the TiB2 content increased from 5 to 10 wt.%. A reduction in penetration depth and horizontal distance were observed as the TiB2 content changed from 10

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to 15 wt.%. However, these changes were not as dramatic as those in the change from 5 to 10 wt.%. The effect of extending the CM time was obvious for the Al5TiB2 samples: a significant depth and distance reduction occurred in the Al5TiB2-20 sample compared to the Al5TiB2-10 sample. The main reason for this result is the increase in surface area of the particles in accordance with the reduced average crystallite size. In other words, the number of particle surfaces that resist applied forces/friction increased with prolonged CM time. The origination of the hard and resistant Al2O3 phase could be another reason for the smaller wear

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ACCEPTED MANUSCRIPT profile, even if only 5 wt.% TiB2 is used for reinforcing the Al matrix. As expected, the lowest changes for the penetration depth and horizontal distance were obtained for the Al15TiB2-20 sample. The origination of Al2O3, especially in the matrix/reinforcement interfaces, may inhibit the ball motion along the sliding direction during the wear tests. An

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increase in corrosion and wear resistance with an increase in TiB2 content was previously reported for TiB2-reinforced AMCs fabricated via stir casting and forging [32]. In addition, the same authors stated that protective Al oxide layers formed at the surfaces of specimens

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depending on the TiB2-content-related microstructure and dispersion, which prevent the underlying surfaces from further corrosion [32]. Considering the formation of the Al2O3 phase

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in the microstructure as a discontinuous layer, it is therefore reasonable that in present study an increase in wear resistance was observed with an increase in the TiB2 content. The wear characteristics of the sintered samples are given in Table 2 in terms of wear volume losses (mm3) and relative wear resistances. The Al5TiB2-10 sample had the highest wear

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volume loss of 1.265 mm3 and the lowest wear resistance of 1.00. The relative wear resistance of each sample was calculated by dividing the wear volume loss of this Al5TiB2-10 sample to that of the measured composite. Consistent with the changes in penetration depth and

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horizontal distance in Fig. S1 (supplementary data), the relative wear resistance increased as the TiB2 content increased at a fixed CM time or as the CM time increased at a constant TiB2

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content. A sharp decrease in the relative wear resistance was observed from the Al5TiB2-10 (resistance 1) to the Al5TiB2-20 (resistance 3.25) sample. The difference in relative wear resistance between the Al10TiB2-10 (resistance 3.54) and Al10TiB2-20 (resistance 3.88) samples, however, was unremarkable. This difference became even smaller as the TiB2 content increased to 15 wt.% for the 10 (resistance 4.58) and 20 min (resistance 4.62) CM’d and sintered samples. Therefore, there may be no notable effect of CM time on the relative wear resistance for Al-10TiB2 and Al-15TiB2 samples. However, amongst all samples, the

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ACCEPTED MANUSCRIPT Al15TiB2-20 sample had the highest wear resistance value. For all specimens sintered from powders that were CM’d for 20 min, no significant difference in relative wear resistance occurred owing to the contribution of the Al2O3 phase to the sliding wear mechanism. In a recent study, in which the effects of Al2O3 particles on AMCs prepared via magneto-milling

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and hot pressing were investigated, it was reported that hardness, strength, and wear resistance of the composites increased with increasing volume fraction of Al2O3 up to 10 vol.% [55]. The high wear resistance caused by the Al2O3 reinforcement particles was

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explained in another study as resistance to abrasion and restriction of the deformation of AMCs [56]. Similarly, Al2O3-reinforced (up to 2 wt.%) AMCs fabricated by the liquid

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metallurgy method coupled with powder metallurgy was reported to exhibit significantly improved mechanical properties [57]. For these reasons, the addition of 15 wt.% TiB2 and the origination of Al2O3 in the Al15TiB2-20 sample resulted in a higher wear resistance value than for the other specimens. It should also be noted that the CM time was clearly related to

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the lowest TiB2 content.

Representative SEM images of the wear tracks taken from the Al5TiB2 and Al15TiB2 samples sintered from the 10-min CM’d powders are shown in Figs. S2(a) and (b) in the

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supplementary data. The Al5TiB2 sample exhibited a wider wear track, in which deep and continuous grooves along the sliding direction can be clearly seen in Fig. S2(a)

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(supplementary data). The Al15TiB2 sample, on the other hand, exhibited a narrower wear track with a flatter surface in the presence of various ruptured regions. Similarly, some surface defects, such as matrix tearing, matrix smearing, voids, microgrooves, scratches, and pits, were formed during the cutting process of TiB2-reinforced AMCs prepared via a mixedsalts method [34, 38]. Based on Figs. S1 and S2 (supplementary data) and Table 2, the changes in penetration depth and horizontal distance, SEM images of wear tracks, and relative wear resistances are all in good agreement between each other. The Al5TiB2 sintered from

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ACCEPTED MANUSCRIPT powders MA’d for 4 h had the lowest wear resistance value, whereas Al15TiB2 had the highest [14]. This result means that an increase in the amount of TiB2 particles dispersed throughout the matrix creates more obstacles, limits the motion of dislocations, hinders the flow of the material, and resists plastic deformation during sliding under applied loads [14,

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24–25, 34]. In a similar way, the tendency of inhibited recrystallization and grain growth in the presence of reinforcement particles and/or some impurities has been reported in some studies [33, 37]. Moreover, the absence of Al3Ti in composites, thanks to the nondeformable

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TiB2 particles due to CM prior to sintering, was the main reason for the high wear resistance values recorded in Al15TiB2 samples [58].

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In addition to the wear characteristics of the composites, the mechanical properties of both the particulate reinforcement and the matrix were determined independently by performing nanoindentation tests. The load–displacement curves obtained from the Al matrix and TiB2 reinforcement of the sintered samples are presented in Figs. 7 and 8, respectively. Indentation

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tests could not be applied to the TiB2 reinforcement for the Al5TiB2-20 sample owing to its low density (95.82%). At the 50-mN maximum force in Fig. 7, the maximal and minimal penetration depths in the matrix were approximately 1860 and 1525 nm, respectively, for

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Al5TiB2-20 and Al15TiB2-20 samples. The maximal and minimal penetration depths in the reinforcement, on the other hand, were found to be approximately 365 and 304 nm,

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respectively, for the Al10TiB2-20 and Al15TiB2-10 specimens, as shown in Fig. 8. As seen in Fig. 8, the curves of the sintered samples shift to lower indentation depths as the TiB2 content in the composites increases from 10 to 15 wt.%, which indicates positive effect of TiB2 particles that was already observed for the relative wear resistances (Table 2). As expected, the amount of plastic deformation dissipated during indentation on the Al matrix (Fig. 7) was much higher than in the TiB2 reinforcements (Fig. 8).

15

ACCEPTED MANUSCRIPT The hardness and elastic modulus values for the composite constituents, obtained according to the Oliver–Pharr method, are listed in Table 3 [59]. It is evident that the values for both hardness and elastic modulus of the matrix and reinforcement were the highest among all samples when 15 wt.% TiB2 and CM times of 10 and 20 min were used (Table 3). For the

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same TiB2 content of 15 wt.%, the elastic modulus of the TiB2 reinforcement decreased from 433.7 to 324.5 GPa, and the hardness decreased from 43.53 to 33.95 GPa, as the CM time increased from 10 to 20 min. The elastic modulus of the Al matrix, in contrast, increased from

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109.2 to 143.2 GPa, and the hardness increased from 0.63 to 0.89 GPa, as the CM time increased from 10 to 20 min. These results are reasonable considering that the microstructural

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characterization presented in Fig. 5 showed evidence of Al2O3 formation in the matrix/reinforcement interfaces that resulted in a reduction of TiB2 strengthening (Table 3). The softening and variations of the hardness and elastic modulus values of the TiB2 reinforcement can be attributed to the densification rates of the composites, which decreased

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with increasing CM time (Fig. 6). The decrease in relative densities at increased CM times had a greater effect on the hardness values of TiB2 than on those of the Al matrix. This indicates that the mechanical energy added by CM played an important role in the

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strengthening of the ductile major Al matrix, whereas the densification of the composites considerably affected the strengthening of the hard and discontinuous TiB2 particulates. With

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respect to the TiB2 content, on the other hand, no comparable trend could be observed between the density of the composites (Fig. 6) and the hardness of the reinforcement (Table 3).

In addition, the ultrahigh hardness and elastic modulus (43.53 and 433.7 GPa, respectively) of TiB2 in the Al15TiB2-10 sample can be attributed to the clean interface without the presence of any secondary phases, which was mainly achieved during CM for 10 min [46]. The elastic modulus of Al measured by nanoindentation testing was previously reported as 68 GPa [59],

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ACCEPTED MANUSCRIPT which is much lower than any of the obtained results in the present study listed in Table 3. This improvement can be attributed to the grain size refinement of the Al phase (65 nm of crystallite size) in the composite powders during 20 min of CM (Table 1), which contributed to the strength of the matrix in the final composites. Thus, a possible strengthening

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mechanism of TiB2 particulate-reinforced Al matrix composites can be related to the grain size refinement of the Al matrix [60]. The rule of mixtures can be applied to predict the hardness of the composites (assuming isotropic materials) according to equation (1) [60–61]. (1)

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Hc = Hmfm + Hrfr

where Hc, Hm, and Hr are the hardness values for the composite, matrix, and reinforcement,

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and fm and fr are the volume fractions of the matrix and reinforcement, respectively. To assess the validity of the rule of mixtures, the hardness values calculated using equation (1) were compared with the experimental results for the 10-min CM’d and sintered samples [61]. Table 4 presents the average Vickers microhardness measurements and the calculated hardness

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values for the selected sintered samples. For the Al5TiB2-10, Al10TiB2-10, and Al15TiB2-10 composites, microhardness values of 0.75, 1.23, and 1.35 GPa were obtained, respectively. Nanoindentation test results of the Al15TiB2-10 sample showed that the hardness values for

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Al and TiB2 were 0.63 and 43.53 GPa, respectively. Taking the data in Table 3 for Hm and Hr, equation (1) gives the hardness of the composite as 1.49 GPa. Similar calculations give Hc

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values of 0.83 and 1.34 GPa for the Al5TiB2-10 and Al10TiB2-10 composites, respectively. The calculated hardness values are summarized in Table 4; they agree well with the experimental values, which corroborates the validity of the utilized model. The absence of any intermediate phase between the matrix and the reinforcement indicates the feasibility of the calculations. Furthermore, both the experimental and theoretical hardness values of the composites were considerably improved over those for monolithic Al and other AMCs reinforced with various particulates [31, 46].

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ACCEPTED MANUSCRIPT For each type of composite, three tensile specimens were tested at room temperature. Fig. 9 shows typical stress–strain curves for the sintered samples obtained from the tensile tests. As seen in Fig. 9, the stress–strain curve for the Al15TiB2-20 composite begins to deviate from the linear elastic region at a stress as low as 50 MPa, although the tensile stress increases

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continuously to approximately 185 MPa (with an elongation up to 2.3 %) after that. This is actually the fracture strength, because the sample fractured before reaching the maximal stress in the stress–strain curve for a completely dense sample [47]. Similar behavior is observed for

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all composites. This is an indication of a problem related to the fabrication of a tensile specimen whose sinterability could not be completely achieved by the utilized cold pressing

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and pressureless sintering method. Tensile specimens of the Al-TiB2 composites have been earlier produced by hot extrusion or friction stir processes to achieve the elongation of coarsegrained regions in the plastic deformation direction [36, 48]. These processes result in a good combination of high strength and remarkable plasticity [15]. On the other hand, the highest

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tensile strength of approximately 171 MPa with a low elongation (up to 0.4%) was obtained for the Al-TiB2 composites fabricated by accumulative roll bonding [40]. Furthermore, ultimate tensile strength of approximately 209 MPa with an elongation of 4.6% was reported

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for the Al-7Si/10TiB2 composites fabricated by the salt reaction route [25]. In compliance with the previous studies, the fracture strengths gradually increased with increasing TiB2

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content, as seen in Fig. 9 [25, 40]. Thus, considering the morphology of the samples revealed in their microstructures, interfacial relations, densities, and wear resistance and hardness values, the microstructural and mechanical characterizations of the final products showed that short CM (10 min) as a powder preparation method prior to pressureless sintering and a high content of TiB2 particles (15 wt.%) enabled the production of sintered Al-TiB2 samples with improved properties. For instance, an increase in the amount of TiB2 particulate reinforcement enabled a decrease in the

18

ACCEPTED MANUSCRIPT average crystallite size of the Al phase, an increase in densification (up to 10 wt.% TiB2), a decrease in penetration depth and horizontal distance changes and hence an increase in the relative wear resistance values of the composites. Thus, the hard and brittle TiB2 increasing in the Al matrix in conjunction with the CM process contributed to the microstructural (average

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crystallite size, morphology) and mechanical properties (hardness, relative wear resistance, elastic modulus, fracture strength) of the composites. TiB2 particulate-reinforced AMCs fabricated in the present study may therefore be potentially used in structural engineering

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components in the aerospace, military, and automotive industries after long-term and realtime performance tests. Overall, considering the simplicity of the current processing route and

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its other advantages—i.e., the reduced milling times and absence of any undesired intermetallic phases—it is expected that this approach will give rise to new investigations into the fabrication of metal boride particulate-reinforced AMCs. 4. Conclusions

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In this study, the effect of cryomilling on the microstructural and mechanical properties of pressurelessly sintered Al-xTiB2 (x = 5, 10, or 15 wt.%) composites was investigated. Based on the results, the following conclusions can be drawn:

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All cryomilled powders contained only Al and TiB2 phases without the formation of any impurities and secondary phases. Cryomilling for 20 min compared to 10 min resulted in a greater reduction in the average crystallite size of the powders. After

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•

sintering of the specimens cryomilled for 20 min, but not for 10 min, the formation of a small amount of the Al2O3 secondary phase in the matrix/reinforcement interface

was observed. The formation of Al3Ti, which was usually observed at the interfaces of the sintered Al-TiB2 composites, was prevented by milling under cryogenic conditions, even when up to 15 wt.% of TiB2 was used. Microstructural evaluations showed a uniform distribution of TiB2 in the Al matrix.

19

ACCEPTED MANUSCRIPT •

Samples cryomilled for 10 min and subsequently sintered had higher densities than those cryomilled for 20 min and sintered, owing to the presence of the Al2O3 phase weakly bonded to the Al matrix.

•

The Al5TiB2-10 sample had the lowest wear resistance of 1.00, whereas the Al15TiB2-

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20 sample had the highest wear resistance of 4.62 owing to the presence of both TiB2 and Al2O3 particles that acted as obstacles and restricted the motion of the sliding ball. Furthermore, the highest fracture strength of approximately 185 MPa with an

•

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elongation of 2.3% was achieved for the Al15TiB2-20 composite.

The elastic modulus and hardness values of the Al matrix and TiB2 reinforcement, and

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their dependence on the varying cryomilling times and TiB2 contents of the composites, were determined separately. The absence of any intermediate phase between the matrix and reinforcement provided the option to calculate the hardness values for the composites. Both the experimental and theoretical hardness values were

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significantly improved compared to those for monolithic Al and Al matrix composites reinforced with other particulates. •

The utilization of 10 min of cryomilling and incorporation of 15 wt.% TiB2 particles

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into the Al-TiB2 composites was found to be more effective than a longer milling time

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or lower TiB2 content. In addition, it was shown that milling under cryogenic conditions prior to pressureless sintering resulted in the absence of any intermetallic phase at the interface between the matrix and reinforcement of the Al-TiB2

composites.

Acknowledgements In this study, the main infrastructure of PML laboratories was used for powder preparation and sintering experiments. The authors acknowledge Prof. Dr. Hüseyin Çimenoğlu and Faiz Muhaffel, M.Sc., for their help with the wear tests. The authors would like to thank Koç

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ACCEPTED MANUSCRIPT University Surface Science and Technology Center (KUYTAM) for instrumental support and Dr. Barış Yağcı and Dr. Amir Motallebzadeh for their help with the characterization studies. Appendix A. Supplementary data The following is the supplementary data related to this article:

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*Supplementary Information File Word Document References

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ACCEPTED MANUSCRIPT Figure 1. X-ray diffraction (XRD) patterns of AlxTiB2-y (x = 5, 10, or 15 wt.%) powders cryomilled for different time periods (y = 10 or 20 min): (a) Al5TiB2-10, (b) Al10TiB2-10, (c) Al15TiB2-10, (d) Al5TiB2-20, (e) Al10TiB2-20, and (f) Al15TiB2-20. Figure 2. XRD patterns of the bulk samples sintered from the AlxTiB2-y powders: (a)

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Al5TiB2-10, (b) Al10TiB2-10, (c) Al15TiB2-10, (d) Al5TiB2-20, (e) Al10TiB2-20, and (f) Al15TiB2-20.

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Figure 3. Backscattered electron scanning electron microscope (SEM) images of the bulk samples sintered from the AlxTiB2-y powders: (a) Al5TiB2-10, (b) Al10TiB2-10, (c)

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Al15TiB2-10, (d) Al5TiB2-20, (e) Al10TiB2-20, and (f) Al15TiB2-20.

Figure 4. Secondary electron SEM image and energy-dispersive x-ray spectroscopy (EDX) maps of the matrix-reinforcement interface obtained for the sintered Al10TiB2-10 sample. Figure 5. Secondary electron SEM image and EDX point analyses of the matrix-

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reinforcement interface obtained for the sintered Al15TiB2-20 sample. Figure 6. Relative densities of the bulk AlxTiB2-y sintered samples and those of

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mechanically alloyed powders (adapted from [14]): sintered samples from powders (a) mechanically alloyed for 8 h [14], (b) mechanically alloyed for 2 h [14], (c) cryomilled for 20

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min, and (d) cryomilled for 10 min.

Figure 7. Load-displacement curves of the sintered samples obtained for the Al matrix. Figure 8. Load-displacement curves of the sintered samples obtained for the TiB2 reinforcement. Figure 9. Stress-strain curves of the sintered samples.

ACCEPTED MANUSCRIPT Table 1. Average crystallite sizes of the Al phase in the cryomilled Al5TiB2, Al10TiB2 and Al15TiB2 powders. Milling Time (min)

Average Crystalite Size (nm) 168

10

113

Al-5TiB2 Al-10TiB2 Al-15TiB2

105

Al-5TiB2

104 20

73

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Al-10TiB2

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Sample Name

65

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Al-15TiB2

ACCEPTED MANUSCRIPT Table 2. Wear characteristics of the sintered samples. Relative Wear Resistance (%) 1

Al-10TiB2-10

0.357

3.54

Al-15TiB2-10

0.276

4.59

Al-5TiB2 -20

0.389

3.25

Al-10TiB2-20

0.326

3.87

Al-15TiB2-20

0.274

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Al-5TiB2-10

Wear Volume Loss (mm3) 1.265

4.62

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Sample Name

ACCEPTED MANUSCRIPT Table 3. Hardness and elastic modulus of the matrix and reinforcement obtained from the nano-indentation tests of the sintered samples. Al matrix

TiB2 reinforcement

Al5TiB2-10

0.63

106.0

30.94

Al10TiB2-10

0.68

107.4

31.13

Al15TiB2-10

0.63

109.2

43.53

Al5TiB2 -20

0.60

109.1

Al10TiB2-20

0.70

101.0

Al15TiB2-20

0.89

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283.0

362.7 433.7

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(GPa)

-

-

27.61

280.7

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Sample Name

Elastic Modulus (GPa)

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Hardness

(GPa)

Elastic Modulus (GPa)

Hardness

33.95

324.5

ACCEPTED MANUSCRIPT Table 4. Average microhardness (Vickers) and the calculated hardness values of the sintered samples. Hardness calculated theoretically (GPa) 0.83

Al-10TiB2-10

1.23 ± 0.17

1.34

Al-15TiB2-10

1.35 ± 0.20

1.49

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Al-5TiB2-10

Average Hardness experimental (GPa) 0.75 ± 0.11

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ACCEPTED MANUSCRIPT Highlights •

Microstructures of Al-TiB2 composites were investigated with respect to their TiB2 content. Cryomilling had positive effects on the uniform distribution of TiB2 particles.

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No formation of intermetallic phases, such as Al3Ti, was observed at the interfaces.

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10-min cryomilling before sintering led to elastic moduli of 283-433.7 GPa (TiB2).

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The absence of interfacial reactions improved the mechanical properties.

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