Effect of milling type on the microstructural and mechanical properties of W-Ni-ZrC-Y2O3 composites

Ceramics International 43 (2017) 7106–7114

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Effect of milling type on the microstructural and mechanical properties of W-Ni-ZrC-Y2O3 composites

MARK

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Duygu Ağaoğullarıa, , Özge Balcıb, M. Lütfi Öveçoğlua a Istanbul Technical University, Faculty of Chemical and Metallurgical Engineering, Department of Metallurgical and Materials Engineering, Particulate Materials Laboratories (PML), 34469 Maslak, Istanbul, Turkey b Koç University, Department of Chemistry, Rumelifeneri Yolu, 34450 Sarıyer, Istanbul, Turkey

A R T I C L E I N F O

A BS T RAC T

Keywords: Milling Sintering Composites Microstructure Mechanical properties

This study reports the effect of milling type on the microstructural, physical and mechanical properties of the WNi-ZrC-Y2O3 composites. Powder blends having the composition of W-1 wt% Ni-2 wt% ZrC-1 wt% Y2O3 were milled at room temperature for 12 h using a Spex™ 8000D Mixer/Mill or cryomilled in the presence of externally circulated liquid nitrogen for 10 min using a Spex™ 6870 Freezer/Mill or sequentially milled at room temperature and cryogenic condition. Then, powders were compacted in a hydraulic press under a uniaxial pressure of 400 MPa and green bodies were sintered at 1400 °C for 1 h under Ar/H2 atmosphere. Phase and microstructural characterization of the milled powders and sintered samples were performed using X-ray diffractometer (XRD), TOPAS software, scanning electron microscope/energy dispersive spectrometer (SEM/ EDS), X-ray fluorescence (XRF) spectrometer and particle size analyzer (PSA). Archimedes density and Vickers microhardness measurements, and sliding wear tests were also conducted on the sintered samples. The results showed that sequential milling enables the lowest average particle size (214.90 nm) and it is effective in inhibiting W grain coarsening during sintering. The cryomilled and sintered composite yielded a lower hardness value (5.80 ± 0.23 GPa) and higher wear volume loss value (149.42 µm3) than that of the sintered sample after room temperature milling (6.66 ± 0.39 GPa; 102.50 µm3). However, the sequentially milled and sintered sample had the highest relative density and microhardness values of 95.09% and 7.16 ± 0.59 GPa and the lowest wear volume loss value of 66.0 µm3.

1. Introduction

provides a short circuit diffusion route in regard of quick mass transportation, reduces the sintering temperature and enhances the densification rate [5–9]. Moreover, some ceramic reinforcing agents such as carbides (TiC, ZrC, HfC, SiC, etc.) and oxides (La2O3, Y2O3, ThO2, ZrO2, etc.) have been used to improve the mechanical properties of W matrix composites [10–16]. On the other hand, the effects of some oxide (La2O3, Y2O3 and Al2O3) and/or boride (TiB2, CrB2, HfB2 and ZrB2) ceramic reinforcements on the properties of Ni activated sintered W matrix composites have been studied in the recent literature [17–21]. Among the ceramic reinforcement particles, zirconium carbide (ZrC) is an important ultra-high temperature ceramic with excellent properties such as high melting point, high temperature strength and good chemical resistance [11,22]. Therefore, there are several investigations reported on the microstructural, physical, mechanical and ablation properties of W-ZrC composites [11,12,22–29]. However, the combined addition of carbide and oxide ceramic reinforcements in the Ni activated W composites has not been studied thoroughly yet.

Metal matrix composites (MMCs) have been used in many industrial applications such as aerospace, automotive, aircraft, electronics especially due to their improved mechanical properties at ambient and elevated temperatures [1]. Amongst MMCs, tungsten (W) matrix composites are candidate materials for important structural applications at high temperatures due to their excellent properties such as high elastic modulus, high thermal shock resistance, low thermal expansion coefficient, good corrosion resistance and good high temperature strength and stiffness [1–3]. However, the densification of W and W matrix composites is difficult since it requires high sintering temperature and long sintering duration due to the high melting point and low ductility of W [2–5]. In order to obtain highly densified W bodies, activated sintering of W and W matrix composites has been utilizing and attracting great attention in the current literature [5–8]. Activated sintering method includes the addition of small amounts of metallic additives (Ni, Co, Fe, Pt, Pd, etc.) into the matrix and it

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Corresponding author. E-mail addresses: [email protected] (D. Ağaoğulları), [email protected] (Ö. Balcı), [email protected] (M.L. Öveçoğlu).

http://dx.doi.org/10.1016/j.ceramint.2017.02.142 Received 30 December 2016; Received in revised form 12 February 2017; Accepted 26 February 2017 Available online 28 February 2017 0272-8842/ © 2017 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

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Fig. 1(e)–(h), the average particle sizes of the W, Ni, ZrC and Y2O3 powders are as 2.52 µm, 2.67 µm, 285 nm and 590 nm, respectively. The content of the Ni activator was selected as 1 wt% of the W matrix. In regard of the W-Ni binary phase diagram, 1 wt% Ni provides a high solubility ratio, does not result in the emergence of a line compound and directly goes into the W solid solution [33]. Also, previous studies carried out in our laboratory facilities revealed that 1 wt% Ni is an adequate agent amount for the activated sintering of W [17–21]. 1 wt% Ni powders were added into the W by mechanical alloying (MA) process which enables a rapid solid-state incorporation of the sintering aid into the matrix with an enhanced activation as well as particle size reduction. Preliminary experiments conducted for the determination of MA duration (up to 18 h) showed that there is not a considerable change in the W crystallite size after 6 h of MA. Thus, 1 wt % Ni added W powders were mechanically alloyed (MA’d) for 6 h and they are hereafter referred to as W1Ni pre-alloy. In order to compensate the particle size difference between the W1Ni pre-alloy and the reinforcement/dispersoid materials as desired in the production of MMCs, ZrC and Y2O3 powders were also pre-milled for 6 h. X-ray diffraction (XRD) investigations of the pre-alloyed W1Ni and premilled ZrC and Y2O3 powders represented in Fig. 2(a) through (c) were carried out using a Bruker™ D8 Advanced Series powder diffractometer (40 kV, 40 mA) with CuKα (λ=0.154 nm) radiation in the 2θ range of 20–90° incremented at a step size of 0.02° at a rate of 2°/min. The International Center for Diffraction Data® (ICDD) powder diffraction files were utilized for the identification of crystalline phases. As seen from Fig. 2(a), only W (ICDD Card No: 04-0806, Bravais lattice: body-centered cubic, a=b=c=0.316 nm) phase is detected in the XRD pattern of the pre-alloyed W1Ni powders due to the very low weight amount of Ni (1 wt%) under the detection limit of the utilized diffractometer ( > 2 wt% of the overall sample). XRD patterns of the pre-milled ZrC and Y2O3 powders in Fig. 2(b) and (c) show intense crystalline diffraction peaks belonging only to the pure material (ZrC, ICDD Card No: 35-0784, Bravais lattice: face-centered cubic, a=b=c=0.469 nm; Y2O3, ICDD Card No: 44-0399, Bravais lattice: base-centered monoclinic, a=1.389 nm, b=0.349 nm, c=0.861 nm) with no trace of impurities. According to the PSA images of the prealloyed W1Ni, and pre-milled ZrC and Y2O3 powders displayed in Fig. 2(d)–(f), their average particle sizes are determined as 283.1 nm, 197.2 nm and 428 nm, respectively. It should be also noted that the pre-alloyed W1Ni have a larger particle size distribution curve in the range of 100–1000 nm than those of pre-milled ZrC and Y2O3. W1Ni pre-alloy, pre-milled ZrC and pre-milled Y2O3 powders were blended in a WAB™ T2C Turbula blender for 1 h to constitute the composition of W1Ni-2 wt% ZrC-1 wt% Y2O3 (7 g at each run; referred to as-blended W1Ni-2ZrC-1Y2O3). The rationale for selecting the fixed amounts of ZrC (2 wt%) and Y2O3 (1 wt%) is based on our previous investigations which indicated that their higher contents in the W matrix resulted in

Furthermore, there are only few investigations regarding the effects of carbides on the properties of W composites developed by conventional powder metallurgy processes [28,30]. Generally, mechanical alloying (MA) at room temperature has been used as powder preparation technique of W-based powders [17–21]. MA is a novel and simple room temperature process which provides the advantages of easy to handle precursor, simple process control, less equipment requirement and low cost [31]. It is based on the repeated fracturing and welding mechanism of the powder particles and depends on the process parameters of mill type, milling speed, time, atmosphere and media, ball-to-powder weight ratio, process control agent, etc [31]. Furthermore, milling of W-based powders at cryogenic or sequential room temperature/cryogenic conditions has not been investigated yet. Cryomilling process which has been commonly used for ductile metal systems such as Al, Zn, Cu, etc. is beneficial for inhibiting the oxidation/nitridation of the powders, stabilizing the thermal and chemical parameters at low temperature and nitrogen atmosphere during the milling and decreasing coarsening tendency of the particles [32]. Thus, the aim of the present study is to report the effects of different milling types (room temperature milling and/or cryomilling) on the microstructure and properties (density, microhardness and wear volume loss) of the W-Ni-ZrC-Y2O3 composites. This study will contribute to the archival literature with the first results of the ZrC and Y2O3 ceramic particulate reinforced and Ni activated sintered W matrix composites prepared via novel powder metallurgy techniques. 2. Experimental procedure Elemental tungsten (W, Eurotungstene™, 99.9% purity) and nickel (Ni, Alfa Aesar™, 99.9% purity) powders were used as the matrix metal and as the activated sintering agent, respectively. Zirconium carbide (ZrC, Alfa Aesar™, 99.5% purity) and yttrium oxide (Y2O3, Alfa Aesar™, 99.99% purity) powders were utilized as reinforcement and dispersoid materials. In order to determine the initial morphologies of the raw materials, W, Ni, ZrC and Y2O3 powders were examined using a JEOL™ JCM-6000Plus NeoScope scanning electron microscope (SEM) operated at 5 kV (Fig. 1(a)–(d)). The specimens for SEM were prepared by suspending them in ethanol (C2H5OH, Fisher Scientific™, 99.5% purity), transferring onto a base plate, drying under air and coating with a thin layer of gold. The average particle sizes of the initial powders were determined using a Microtrac™ Nano-flex particle size analyzer (PSA) equipped with a Bandelin Sonopuls™ ultrasonic homogenizer using distilled water as the aqueous media. The particle size measurement is essential to represent the larger sizes of the matrix particles than those of reinforcement/dispersoid, which is compatible with the general idea for the preparation of MMCs. According to the particle size distribution curves of the raw materials illustrated in

Fig. 1. SEM and PSA images of the raw materials: (a, e) W, (b,f) Ni, (c,g) ZrC and (d, h) Y2O3.

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Fig. 2. XRD patterns and PSA images of the pre-alloyed and pre-milled powders: (a, d) W1Ni, (b,e) ZrC and (c, f) Y2O3.

Scientific™ XL3t Niton X-ray fluorescence (XRF) spectrometer. Densities of the sintered composites were determined by Archimedes method using ethanol and the result of each sample was reported as the arithmetic mean of five different measurements. After applying a series of metallographic treatments, microstructures of the sintered samples were examined using a Hitachi™ TM-1000 SEM. The polished surfaces was chemically etched using a hot aqueous HF/HNO3 solution with a molar ratio of 4:1 [35]. Microstructures of the etched samples were characterized by utilizing a FEI™ Quanta FEG 250 SEM operated at 10 kV. Vickers microhardness measurements of the sintered samples were conducted using a Shimadzu™ HMV Microhardness Tester under a load of 100 g (9.807 N) for 10 s and the result of each sample includes the arithmetic mean of twenty successive indentations with a standard deviation. Sintered samples were subjected to sliding wear tests at room temperature in a laboratory atmosphere using a Tribotech™ Oscillating Tribotester with a 6 mm alumina ball under an applied force of 4 N, a sliding speed of 4 mm/s and a stroke length of 2 mm for a total sliding distance of 20000 mm. Wear tracks were examined by using a Veeco™ Dektak 6M Stylus profilometer and were imaged using a Leica™ ICC50 HD optical microscope (OM). Wear test result of each sample includes the arithmetic mean of three different profilometer measurements.

lower relative density and microhardness values [18–21]. Room temperature milling was carried out in a Spex™ 8000D Mixer/Mill (at a rate of 1200 rpm) for 12 h using a tungsten carbide (WC)-cobalt (Co) vial (50 ml capacity) and WC-Co balls (ϕ 6.35 mm) with a ball-to-powder weight ratio (BPR) of 7:1. Pre-alloying of W1Ni and pre-milling of ZrC and Y2O3 were also conducted at the same conditions with those used in room temperature milling experiments. Cryomilling was conducted in a Spex™ 6870 Freezer/Mill (at a rate of 900 collisions/min) using a cylindrical polycarbonate vial and stainless steel rods. Cryogenic condition was provided with liquid nitrogen (Linde™, refrigerated) externally circulated around the milling vial. Cryomilling duration was selected as 10 min since preliminary experiments continued up to 20 min resulted in a high amount of wear in the milling rods. Loading, sealing and unloading of the milling vials were done under Ar gas (Linde™, 99.999% purity) in a Plaslabs™ glove box. Hereafter, as-blended W1Ni-2ZrC-1Y2O3 powder milled for 12 h at room temperature is referred to as S1, cryomilled for 10 min in the presence of externally circulated liquid nitrogen is referred to as S2 and sequentially milled at room temperature for 12 h and cryogenic condition for 10 min is referred to as S3. Then, S1, S2 and S3 powders were compacted in a 10 t capacity MSE™ MP-0710 uni-action hydraulic press to obtain cylindrical preforms with a diameter of 6.5 mm under a uniaxial pressure of 400 MPa. The green bodies were sintered at 1400 °C for 1 h in a Linn™ HT-1800 high-temperature controlledatmosphere furnace with a heating and cooling rate of 10 °C/min under Ar/H2 gas flowing conditions. XRD analyses of the as-blended and milled powders, and sintered composites were carried out using the same diffractometer with the same conditions utilized for the pre-alloyed and pre-milled powders. The average crystallite sizes and lattice strains of the W phase in the milled powders were determined using a Bruker™-AXS TOPAS 4.2 software using the modified Scherrer's formula based on the broadening of the most intense XRD diffraction peaks (with (110), (200) and (211) reflections) fitted according to the Lorentzian profile by applying fundamental parameters approach [34]. Particle size measurements of the milled powders were conducted using the same method utilized for the pre-alloyed and pre-milled powders. Energy dispersive spectroscopy (EDS) analysis of the as-blended powders was performed using the same SEM (operated at 5 kV) utilized for the raw materials. The amounts of Co contamination in the as-blended powders and powders milled at room temperature for 12 h were detected by using a Thermo

3. Results and discussion 3.1. Phase and microstructural analyses of the as-blended and milled powders Prior to the room temperature milling and cryomilling processes, it is essential to show the structure of the as-blended W1Ni-2ZrC-1Y2O3 powders containing pre-alloyed W1Ni, pre-milled ZrC and pre-milled Y2O3 powders. Fig. 3 illustrates the XRD pattern of the as-blended W1Ni-2ZrC-1Y2O3 powders. Only W and ZrC phases are detected in the XRD pattern of the as-blended powders since the weight amounts of Ni and Y2O3 are under the detection limit of the diffractometer. The SEM image of the as-blended powders in Fig. 4(a) shows an agglomerated microstructure even though the pre-alloyed W1Ni and the premilled ZrC and Y2O3 have nano-sized particles. Fig. 4(b)–(g) illustrate the corresponding EDS mappings of the SEM image in Fig. 4(a), showing the elemental maps for W, Ni, Zr, C, Y and O, respectively. The W elemental map (Fig. 4(b)) coincides almost entirely with the Ni 7108

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expected since they were not observed in the XRD pattern of the asblended powders (Fig. 3). However, the absence of the ZrC phase is probably due to its peak broadening caused by the continuous deformation during milling. The repeated welding/fracturing mechanism during milling processes can result in decomposition/degradation of the particles and this can trigger the formation of secondary phases [31]. However, any diffraction peak belonging to a secondary phase could not be detected after different milling types. Furthermore, XRD patterns of the S1 and S3 samples do not show the emergence of WC contamination worn from the WC-Co vial and balls during 12 h of milling. The absence of WC contamination can be also explained with its very small content in the W matrix. For that reason, XRF analyses were conducted on the S1 sample due to detect the Co contamination worn/torn from the milling vial/balls. It was found that 12 h of room temperature milling causes a Co contamination of about 0.38 wt%, proving also the presence of WC contamination under the detection limit of diffractometer. In order to understand the effect of different milling types at room temperature and/or cryogenic condition on the as-blended W1Ni2ZrC-1Y2O3 powders, the average crystallite sizes and lattice strains of the W phase in the milled powders (S1, S2 and S3), and average particle sizes of these milled powders were determined (Table 1). Taking into account the initial average particle sizes of the pre-alloyed W1Ni and pre-milled ZrC and Y2O3, it can be stated that milling of the as-blended W1Ni-2ZrC-1Y2O3 powders at room temperature for 12 h resulted in a significant decrease in the average particle size of the S1 powders (290.50 nm). The average crystallite size and lattice strain of the W phase in the S1 sample were found as 8.0 nm and 3.166%. This decrease is a result of continuous fracturing and welding mechanism triggered by the high impact energy released from the collisions of powders with vial and balls during milling at room temperature. In the case of cryomilling of the as-blended W1Ni-2ZrC-1Y2O3 powders for 10 min (S2), the average crystallite size of the W phase is about 1.76 times higher (14.10 nm) and the lattice strain is 1.8 times lower (1.755%) than those of the S1 sample. Additionally, the XRD pattern of the S1 sample in Fig. 5(a) also has lower calculated area of W reflections than those of S2. This supports the fact that higher broadening of the W peaks results in higher decrease in the average crystallite size and increase in the lattice strain. Considering the milling time without including the operational differences of the milling processes, it can be said that only cryomilling for 10 min is useful to obtain an approximate refinement with that of milling at room temperature for 12 h. Moreover, the average particle size of the S2

Fig. 3. XRD pattern of the as-blended W1Ni-2ZrC-1Y2O3 powders.

elemental map (Fig. 4(c)), suggesting the presence of W1Ni pre-alloy. The Zr elemental map (Fig. 4(d)) clearly overlaps with the C elemental map (Fig. 4(e)), indicating the presence of ZrC phase. Besides, the occurrence of the Y2O3 phase can be proven by the Y and O elemental maps (Fig. 4(f) and (g)) which fit the same regions in the microstructure. Especially, the elemental maps of Ni, Y and O (Fig. 4(c), (f) and (g)) are useful as the proofs of the Ni sintering agent and the Y2O3 dispersoid which were not detected in the XRD patterns (Figs. 2(a) and 3). It should be noted that the accumulation of the elements in the same regions (Fig. 4(b)–(g)) is an expected character of the as-blended powders in which a well distributed microstructure could not be provided. Moreover, any signal of element other than W, Ni, Zr, C, Y and O was not observed in the as-blended powders during EDS analyses. Considering limited resolution power of elemental mapping, XRF analysis was also employed on the as-blended powders and it confirmed the absence of impurity. Fig. 5(a) illustrates the XRD patterns of the milled W1Ni-2ZrC1Y2O3 powders using different milling types at room temperature and/ or cryogenic condition (S1, S2 and S3). As seen from Fig. 5(a), all milled powders contain only the W phase. No peaks belonging to the Ni, ZrC and Y2O3 phases are identified in the XRD patterns after milling at room temperature for 12 h, cryomilling for 10 min or sequential milling process. The absence of Ni and Y2O3 is already

Fig. 4. SEM/EDS analyses of the as-blended W1Ni-2ZrC-1Y2O3 powders: (a) SEM image, and elemental maps for (b) W, (c) Ni, (d) Zr, (e) C, (f) Y and (g) O.

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Fig. 5. XRD and SEM analyses of the milled W1Ni-2ZrC-1Y2O3 powders using different milling types at room temperature and/or cryogenic condition (S1-milled at room temperature for 12 h, S2-cryomilled in the presence of externally circulated liquid nitrogen for 10 min, and S3-sequentially milled at room temperature for 12 h and cryomilled for 10 min): (a) XRD patterns of S1, S2 and S3, and SEM images of (b) S1, (c) S2 and (d) S3.

cryomilling was applied to the as-blended W1Ni-2ZrC-1Y2O3 powders (S2) instead of room temperature milling for 12 h, the sizes of the agglomerates decreased to about 2.5 µm and below (Fig. 5(c)). Furthermore, when sequential milling was conducted on the asblended W1Ni-2ZrC-1Y2O3 powders (S3), it is clearly seen that large agglomerates broke down and irregular agglomerates in sizes not larger than 1 µm appeared in the microstructure. Thus, it can be stated that SEM images in Fig. 5(b)–(d) conform well to the average particle size measurements.

sample is lower (248.20 nm) than that of S1, indicating that cryomilling enables the prevention of agglomeration. Although the milling for S1 sample was conducted at room temperature without applying any external heat, there should be a temperature increase inside the vial due to the repeated collisions. The repeated fracturing and welding mechanism in the presence of an amount of temperature increase could result in a higher average particle size value in the S1 sample than that of S2, even if longer milling duration was used. Since any temperature increase did not occur during cryomilling in the presence of externally circulated liquid nitrogen and rewelding mechanism was hindered, S2 sample had a lower average particle size than that of S1. Thus, the combined process of milling at room temperature for 12 h and cryomilling for 10 min should be more beneficial for particle refinement and prevention of agglomeration. The average crystallite size (7.90 nm) and lattice strain (3.181%) values of the W phase in the sequentially milled powders (S3) are close to those of the S1 sample. However, there is a remarkable decrease in the average particle size of S3 (214.90 nm), compared to that of S1 (290.50 nm). This means that subsequent 10 min of cryomilling provided breaking down of the agglomerated particles. Fig. 5(b) through (d) illustrate the SEM images of the S1, S2 and S3 powders. All the powders milled at room temperature and/or cryogenic condition do not consist of perfect spheroidal-shaped particles throughout the structure because repeated collisions during milling create morphologically inhomogeneous particles such as having spherical, irregular and equiaxed-shaped structures in the same batch [31]. It is also easily seen from the microstructures that some smaller roundshaped particles are embedded in large agglomerates. This prevents the observation of accurate sizes for the smaller particles. Amongst Fig. 5(b)–(d), the largest agglomeration (up to about 4 µm) is observed in the microstructure of the S1 powders. When only 10 min of

3.2. Phase and microstructural analyses of the sintered samples Fig. 6 shows the XRD patterns of the bulk samples sintered from the S1, S2 and S3 powders. The XRD peaks of the sintered samples became sharp, narrow and tall in intensity, indicating improved crystallinity due to the effect of high temperature during sintering. All sintered samples exist the peaks of Ni(W) solid solution phase and WC1−x (ICDD Card No: 20-1316, Bravais lattice: face-centered cubic, a=b=c=0.424 nm) in addition to the major W phase which was already present in the milled powders (Fig. 5(a)). It is already known that the addition of small quantities of Ni provides a high solubility ratio and dramatic densification in the W-Ni system whose line compounds such as Ni4W, NiW and NiW2 could only be favored kinetically over 14.1 wt % Ni [33]. Thus, 1 wt% Ni did not result in the emergence of a line compound and directly went into the W solid solution. In a previous study related with the Ni-W sintered compacts fabricated via MA and pressureless sintering of Ni-xW (x=20, 30 and 40 wt%) powders, XRD analysis revealed the presence of the Ni(W) solid solution phase at the same 2θ value [36]. Additionally, no diffraction peaks of the ZrC and Y2O3 are observed after sintering process, probably arising from their small amounts and broad peaks dominated by the intense peaks of the

Table 1. Average crystallite sizes and lattice strains of the W phase in the powders milled using different milling types at room temperature and/or cryogenic condition, and their average particle sizes. Sample name

Milling type

Average crystallite size (nm)

Lattice strain (%)

Average particle size (nm)

S1 S2

Room temperature milling under Ar atmosphere for 12 h Cryomilled in the presence of externally circulated liquid nitrogen for 10 min Sequentially milled at room temperature for 12 h and cryomilled for 10 min

8.00 14.10

3.166 1.755

290.50 248.20

7.90

3.181

214.90

S3

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sample (Fig. 7(b), (e) and (h)). Although the use of different milling processes did not change the type of the formed phases after sintering (Fig. 6), it strongly affected the microstructure of the samples (Fig. 7(a)–(i)). On the basis of the SEM micrographs of the etched surfaces in Fig. 7(g), (h) and (i), it can be inferred that S1 and S3 samples have the distinct grain boundaries below 1 µm whereas S2 sample has grain boundaries below 8 µm. This means that homogenization of the inclusion particles provided by the 12 h of room temperature milling and sequential milling processes is effective in hindering W grain coarsening than that obtained after 10 min of cryomilling. The accumulated energy of the particles during milling by plastic deformation, fracture, friction, etc. is released more easily in the S2 sample while sintering. Thus, the pinning effect on the growth of W grains could be formed by room temperature milling and sequential milling processes. It should be noted that reduction in grain size and formation of high fraction grain boundaries by milling significantly improves the diffusivity, solid solubility [37] and sintering ability and hence contributes to the physical and mechanical properties of the materials.

Fig. 6. XRD patterns of the bulk samples sintered from the S1, S2 and S3 powders.

W phase. Likewise, XRD pattern of the sintered W1Ni-1 wt% x-2 wt% y (x=La2O3 and Y2O3; y=CrB2, HfB2 and ZrB2) samples in a reported study do not reveal the characteristic peaks of the boride and oxide phases [20]. WC contamination is observed in very weak intensity after sintering (Fig. 6), although it could not be detected in the XRD patterns of the milled powders (Fig. 5(a)) because of its very low content in the W matrix. The presence of WC1−x contamination was also reported by Ağaoğulları et al. for the sintered W1Ni-1 wt% x-2 wt% y (x=La2O3 and Y2O3; y=CrB2, HfB2 and ZrB2) samples milled at room temperature for 12 h in the Spex™ 8000D Mixer/Mill using WC-Co vial/balls with a BPR of 7:1 [20]. However, it is surprising to detect a very small amount of WC1−x contamination in the S2 sample sintered from the 10 min of cryomilled powders. This can be related with the decomposition of a small amount of ZrC and hence the interaction of carbon and W matrix during sintering process. Similarly, the decomposition of TiB2 reinforcement and formation of W2B phase was reported after sintering of the TiB2 (2, 3 and 4 wt%) incorporated W1Ni matrix [19]. W5Si3 and W2C intermetallic compounds were formerly identified after sintering of W4 wt% SiC-x wt% Y2O3 (x=1 and 5) samples at 1680 °C for 1 h [14]. Also, the occurrence of W2B and NiTi phases were reported for the sintered (1400 °C for 1 h) W1Ni-0.5 wt% Y2O3-x wt% TiB2 (x=4 and 5) samples [17]. Thus, the emergence of WC contamination is due to probable wear of WC from the vial and balls during room temperature milling for 12 h and/or slight reaction of decomposed ZrC with W matrix. It should be also noted that increasing the content of the reinforcement over 2 wt% in the overall composition could result in the emergence/enhancement of the contamination and intermetallics, on the basis of Fig. 6 and the reported literature [14,17,19]. SEM images taken from the polished (Fig. 7(a)–(c)), fracture (Fig. 7(d)–(f)) and etched (Fig. 7(g)–(i)) surfaces of the bulk samples sintered from the S1, S2 and S3 powders are useful for comparing their microstructural differences resulting from different milling types. As seen from Fig. 7(a) and (c) and Fig. 7(d) and (f), S1 and S3 samples have similar microstructures, indicating the effect of room temperature milling for 12 h. Nano-sized ZrC and Y2O3 particles (dark contrast) which are located both at the grain boundaries and at the grain interiors of W1Ni matrix (light contrast) have a homogenous distribution throughout the polished and fracture surfaces of the S1 and S3 samples. The addition of a second insoluble phase by high energy ball milling such as oxide dispersoid, boride reinforcement, etc. enables smaller W grains by the inhibition of grain growth and retardation of coarsening during the sintering process [20,31]. This phenomena can obviously be supported by the SEM images of the S1 (Fig. 7(a), (d) and (g)) and S3 (Fig. 7(c), (f) and (i)) samples. However, there are large W grains and inhomogeneously distributed and clustered large ZrC and Y2O3 particles in the polished, fracture and etched surfaces of the S2

3.3. Physical and mechanical properties of the sintered samples The density, microhardness and wear volume loss values of the bulk samples sintered from the S1, S2 and S3 powders are given in Table 2. S1 sample has the lowest density value and higher microhardness than that of S2 sample. S3 sample sintered from the sequentially milled powders has the highest relative density and the highest microhardness. The application of cryomilling for 10 min after room temperature milling for 12 h results in an increase in the relative density value (∼6.0%) and in the microhardness value (∼0.5 GPa) of the samples (S1 and S3). The evolution of average crystallite/particle sizes and density/ microhardness values of the S1 and S3 samples are in great consistency with each other, in which the decrease in average crystallite/particle size corresponds to an increase in density/microhardness. Sintered S2 sample has the lowest microhardness value due to the inhomogeneous distribution of ZrC and Y2O3 particles and growth of W grains. The relative density values of the samples are in good correlation with the SEM images (Fig. 7(a)–(i)) because the presence of pores can be easily observed in their microstructures. Different milling processes employed for the samples are not enough for achieving full densification. Actually, the measured relative density value for the S3 sample is adequately high considering the powder metallurgy methods and sintering temperature used in the present study. Similar relative density (∼93–98%) and microhardness values (∼7 GPa) were reported for the W1Ni-1 wt% x-2 wt% y (x=La2O3 and Y2O3; y=CrB2, HfB2 and ZrB2) samples sintered from the powders milled at room temperature for 12 h [20]. Fig. 8(a)–(c) display the OM images taken from the wear tracks of the bulk samples sintered from the S1, S2 and S3 powders. The worn surfaces of the S1 and S3 samples exhibit continuous grooves along the sliding direction (Fig. 8(a) and (c)) whereas S2 sample has a ruptured wear surface (Fig. 8(b)). As compatible with the OM images in Fig. 8(a)–(c), the highest wear volume loss value was found for the S2 sample and the lowest wear volume loss value was obtained for the S3 sample. Besides, S1 sample has a wider groove than that of S3 sample. ZrC and Y2O3 particles homogeneously distributed throughout the matrix restrict the sliding motion of the hardfacing alumina ball during wear tests and resist rupturing and plastic deformation in the sliding surface. As expected, microhardness results of the samples conform well to the wear volume loss values. It is important to emphasize that the wear volume loss values of the S1 and S3 samples are respectively 3.3 and 5.2 times lower than that of W1Ni-2ZrB2-1Y2O3 samples (347.39 µm3) sintered from the powders milled at room temperature for 12 h [20]. It can be stated that milling of W-based powders at cryogenic or sequential room temperature/cryogenic conditions was investigated at

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Fig. 7. SEM images taken from the polished (a, b, c), fracture (d, e, f) and etched (g, h, i) surfaces of the bulk samples sintered from the S1 (a, d, g), S2 (b, e, h) and S3 (c, f, i) powders.

based systems. However, it has been recently reported that cryomilling can be used to obtain nanocrystalline structures of some Mg alloys [38], oxide dispersion strengthened steels [39] and carbon nanotube reinforced Al composites [40] with enhanced thermal stability and better sintering ability which would further mean improved mechanical properties. In an applicative point of view, it should be also mentioned that powder blends can be reduced to nano-sizes in a relatively short time due to the cryogenic temperature [40,41]. The effectiveness of cryomilling on the milling behavior of Y2O3 dispersion strengthened steels has been investigated and the results showed that cryomilling produced much finer particle/grain size than conventional room temperature milling [39]. Furthermore, cryomilling effectively changed the microstructure of oxide dispersion strengthened steels, and it has a significant effect on creep properties [39]. On the other hand, cryomilling as a powder preparation method before sintering has not been

Table 2. Density, microhardness and wear volume loss values of the bulk samples sintered from the S1, S2 and S3 powders. Sample name

S1 S2 S3

Theoretical density (g/ cm3)

Relative Archimedes density (%)

Microhardness (GPa)

Wear volume loss (µm3)

17.92

88.89 94.41 95.09

6.66 ± 0.39 5.80 ± 0.23 7.16 ± 0.59

102.50 149.42 66.00

first time in this study. Considering the technology and properties of the composites, sequential milling is more effective to yield microstructurally and mechanically improved products for the brittle W-

Fig. 8. OM images taken from the wear tracks of the bulk samples sintered from the S1 (a), S2 (b) and S3 (c) powders.

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genic and ambient conditions and related characterization investigations” and with the project number of 110M130. The authors of this study would like to acknowledge to M.Sc. Emre Tekoğlu and M.Sc. Merve Küçük for their help in SEM and PSA analyses and also to M.Sc. Yakup Yürektürk for his help in wear tests.

previously used for Ni activated sintered W matrix composites reinforced with ceramic particulates. However, there has been a study about the preparation of nano-sized WC-Co powders by cryomilling and about the effects of cryomilling on the microstructural and mechanical properties of these powders [41]. The mechanical properties of nano-sized WC-Co systems produced by cryomilling in the presence of liquid nitrogen internally circulated in the vial were not found to be superior to those of other methods. The presence of high nitrogen and oxygen contents and their reactions with carbon during sintering process negatively affected the sintering behavior and the final properties [41]. The type of cryomill (with external nitrogen circulation system) utilized in the present study prevented the direct contact of liquid nitrogen with powder blends and hence the undesired reactions between the components. Consequently, the evaluation of the measurements carried out for the present samples showed that sequential milling at room temperature and cryogenic condition could be a beneficial process to improve the resultant properties (density, microhardness, wear, etc.) of the W-based composites.

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4. Conclusions In this study, the effect of different milling types (milling at room temperature for 12 h, cryomilling for 10 min or sequential milling process at room temperature and cryogenic condition) were investigated on the microstructural, physical and mechanical properties of the W-Ni-ZrC-Y2O3 composites. Based on the results reported in the present study, the following conclusions can be drawn:

•

•

• • • • •

XRD patterns of all milled W1Ni-2ZrC-1Y2O3 powders exhibited only W phase. Amongst all milled powders, the lowest average crystallite size (7.90 nm) and the highest lattice strain (3.181%) values of the W phase and also the lowest average particle size (214.90 nm) were obtained in the sequentially milled powders. The sequential milling at room temperature for 12 h and cryomilling for 10 min is a more effective route in the particle refinement and breaking down of the agglomerated particles. All sintered samples had the peaks of the dominant W and small amounts of Ni(W) solid solution and WC1−x phases. The peaks of WC contamination emerged after sintering arising from the wear of WC-Co from the vial/balls during room temperature milling for 12 h and/or slight decomposition of ZrC due to the effect of high temperature. In the microstructures of the bulk samples sintered from the powders milled at room temperature and milled with sequential process, there were homogeneously distributed submicron-scale ZrC and Y2O3 particles located both at the grain boundaries and at the grain interiors of W1Ni. There are large W grains and inhomogeneously distributed/clustered large ZrC and Y2O3 particles in the microstructure of the bulk sample sintered from the cryomilled powders. Sequential milling was found as the most effective process in blocking of the W grain coarsening and growth by homogenization of the inclusion particles. Cryomilled and sintered sample exhibited inferior hardness value (5.80 ± 0.23 GPa) than that of the sintered sample after room temperature milling (6.66 ± 0.39 GPa). Sequentially milled and sintered sample had the highest relative density (95.09%), the highest microhardness (7.16 ± 0.59 GPa) and the lowest wear volume loss (66.0 µm3) values.

Acknowledgements This study was supported by The Scientific and Technological Research Council of Turkey (TUBITAK) with the project title of “Development of tungsten based hybrid composites via activated sintering and mechanical alloying using high energy milling at cryo7113

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