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Thermal stability and magnetic saturation of annealed nickel–tungsten and tungsten milled powders
Int. Journal of Refractory Metals and Hard Materials 30 (2012) 78–84
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Int. Journal of Refractory Metals and Hard Materials j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / I J R M H M
Thermal stability and magnetic saturation of annealed nickel–tungsten and tungsten milled powders Kasonde Maweja a,⁎, M.J. Phasha b, L.J. Choenyane a a b
Element Six Ltd/Diamond Research Laboratory, PO Box 561, Springs 1559, Gauteng, South Africa The Council for Scientific and Industrial Research, Materials Science and Manufacturing, Pretoria, South Africa
a r t i c l e
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Article history: Received 12 June 2011 Accepted 13 July 2011 Keywords: Nickel Tungsten Mechanical alloying Annealing Magnetic saturation
a b s t r a c t Crystal structures, microstructures and magnetic saturation of annealed pure W powder along with W–40 wt.% Ni powder mixtures processed by high-energy ball milling were investigated using XRD, DTA, SEM and saturation magnetization techniques. Thermally induced transformations occurred at low temperature annealing. Supersaturated metastable Ni(W) solid solution formed during mechanical milling decomposed during annealing treatment into FCC Ni-rich, FCC W-rich phases and an eta-type phase which was constituted of BCC lattice of W enveloped by two FCC lattices of Ni and W. The structures of the major annealing products were close to Ni10W and W3Ni2. The magnetic saturation of the milled W powder and W–Ni mixtures decreased with the increase in annealing temperature. Milling time was more influential on the magnetic properties of the annealed pure W powders. © 2011 Elsevier Ltd. All rights reserved.
1. Introduction Magnetic hard materials tend to find more and more applications in automotive industry (driving engines, clutches and frictionless bearings), in electronic (starters, sensors) and the like. A real breakthrough in the development of magnetic hard materials turned out to be the discovery, in 1984, of the Nd2Fe14B compound, which showed strong magnetocrystalline anisotropy characteristics [1,2]. The annealing temperature exerted significant impacts upon both structures and magnetic properties of the Nd–Fe–W–B magnets; in this case the effect of annealing temperature was attributed to the formation of Fe7W6 compounds at high annealing temperature [3]. Mondal et al. [4] showed that coarsening particles of mechanically alloyed Fe/FeNi powders impaired the magnetic properties, thus attempt of magnetic application of mechanically alloyed powders necessitates appropriate control over the precipitation/coarsening of the soft magnetic phases and temperature dependent magnetic response of the non magnetic matrix. Rabanal reported that the magnetic saturation of Mg-ferrite soft magnetic materials kept nearly constant, whereas the coercivity increased with milling time [5], whereas Jartych et al. [6] reported the decrease of magnetization versus milling time of mechanically alloyed Fe–W caused by the coexistence of the magnetic Fe(W) and the paramagnetic W(Fe). The studies of the magnetic properties showed that the coercivity decreased and saturation magnetization increased due to micro-stress elimination and reduction in surface dead layers and slight increase in
⁎ Corresponding author. Tel.: + 27 83 365 0952(cell). E-mail address: [email protected] (K. Maweja). 0263-4368/$ – see front matter © 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.ijrmhm.2011.07.005
ferromagnetic ordering upon annealing of Ni58Fe12Zr10Hf10B10 nanostructured mechanically alloyed powders [7]. Bahgat et al. [8] reached similar conclusion for W–Ni–Fe alloys prepared by mechanochemical reduction of mixtures corresponding oxide powders. Nickel–tungsten alloys have received attention as environmentally friendly substitute for hard chrome plating [9] and also because of their enhanced catalytic activity [10]. Activated sintering, that involves small additions of transition metals such as Ni, Co and Fe also enables major reductions in the sintering temperature of W [11–14]. XRD patterns of the sintered samples revealed the presence of (Ni,W) solid solution phase and a small peak of elemental Ni, whereas no intermetallic phases of Ni and W were formed during sintering [15]. On the other hand, thermally induced amorphization of the supersaturated solid solutions formed by mechanical alloying of elemental powders to form Zr52.5Cu17.9Ni14.6 Al10Ti5 and Zr70Pd30 products was observed by DSC and XRD analyses by El-Eskandarany et al. [16] and by Siegrist et al. [17] respectively. The amorphization process preceded the formation of crystalline phases upon annealing [17]. The degree of structural disorder induced by ball milling, the composition evolution upon annealing and the nanograin size distributions within the granular mixed ferromagnetic/anti-ferromagnetic (FM/AFM) powders are key factors for the magnetic properties in granular systems [18–20]. It has been reported that the magnetic saturation of Ni-based mechanically alloyed powders with Fe, Mo and Nb slightly increased at the early stage of milling before it decreased as a result of the electronic interactions between magnetic and non-magnetic elements and finally increased by the partial crystallization of the amorphous matrix [21]. By alloying the Ni–Fe alloys with ternary elements, such as Mo and changing the real structure from crystalline to the amorphous state, various kinds of
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magnetic, electrical and structural properties could be achieved [22,23]. Nickel is known to be a magnetic and tough material, whereas tungsten is hard but is not magnetic and has a high melting temperature (~ 3422 °C). Development of new tough magnetic structural materials for high temperature applications are therefore expected by alloying the two metals. The powder metallurgy route would present less technical challenges inherent to the high melting temperature of W for the manufacturing of this type of alloys. Hence, in the current study the effects of milling time and annealing temperature on the thermally induced transformations and magnetic saturation of W and W–Ni powders are reported.
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Vacuum gold sputter and were analyzed in the JEOL JSM-7500F Field Emission SEM equipped with EDX. Magnetic saturation of the annealed powders was determined using a commercial saturation induction measuring system (LDJ Electronics Inc.). The system was calibrated using pure cobalt and pure nickel references. The test piece of the milled powder was weighed and inserted into a non-magnetic holder. It was inserted into the high-intensity permanent magnet (~0.75 T), and then quickly withdrawn. The response of a search coil was displayed on a magnetic multimeter in terms of magnetic moment (saturation). Each measurement was repeated twice to check consistency of results. 3. Results and discussion
2. Material and experiments 3.1. Thermal analysis of milled powders Ball milling experiments of 100 g W powder at ~99.5% purity (1–2 μm) and 100 g powder mixtures of 60 wt.% W with 40 wt.% Ni (~99.5% purity, 2–3 μm) were conducted in a Retsch PM400 MA-type planetary high energy ball mill in an argon atmosphere at a rotation speed of 350 rpm. The grinding media consisted of 1 kg of 5 mm diameter 3Cr12 stainless steel balls. Stearic acid was added as 3 wt.% of total powder mass as process control agent to minimize cold welding. Milling times of 12 h, 24 h and 48 h were employed. The transformation temperatures of milled powders were determined by differential thermal analysis at a 10 °C/min heating rate in DTA-TGA equipment in an argon atmosphere. In order to characterize the thermally induced transformations, milled powders were annealed near the transformation temperatures for 1 h in a furnace filled with argon. The annealed products were then analyzed by XRD and SEM techniques. The annealed powders were cold mounted in epoxy resin and were mechanically polished using a series of SiC grit papers to cut through the cross-section of the powder particles. The metallographic polishing of samples was achieved using 1 μm diamond pastes. The samples were coated with gold using a Denton
The DTA heating curves of milled W powders and those of W–Ni mixtures are shown in Fig. 1(a) and (b) respectively. The peak temperatures of thermally induced transformation of the powders are indicated by the arrows. The non-milled powders of W and that of W–Ni mixtures were thermally stable throughout the temperature range used (25 °C to 1400 °C) as no transformation peak was observed in the respective heating curves. The curves of W powder (Fig. 1a) milled for 12 h had a small peak around 700 °C, whereas two small broad peaks were observed at ~680 °C and at ~730 °C in W powders milled for 24 h and a single sharp strong peak at ~720 °C was observed in the curves of powders milled for 48 h. All these peaks pointed toward exothermic character, which indicate that transformations, and not only strain relief, took place during the heating of milled W powders. On the other hand, the heating curves of W–40 wt.% Ni powder mixtures (Fig. 1b) milled for 12 h had two broad peaks at ~660 °C and 720 °C. The first exothermic peak shifted at ~ 670 °C for powders milled for 24 h and at around 680 °C for the W–Ni powders milled for
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48 h and its intensity was higher with increased milling time. In addition, the two exothermic peaks observed in W–40 wt.% Ni powders are further apart upon increased milling. First one positioned below 700 °C while the second lies above 800 °C as shown by the blue arrows in Fig. 1b. Due to inhomogeneity of ball milling, some fine and deformed W particles would transform earlier/faster than those that get exposed to severe deformations at a later stage of milling. The evolution of the particle sizes of pure W and W–40 wt.% Ni mixtures are shown in Fig. 2. The particle size distribution parameters of milled W powders presented minima at 12 h (Fig. 2a), which indicates that fracturing of the powder particles prevailed at the early milling stage. The d10, d50 and d90 of the powders were smaller than those of the starting powders and those of the products milled longer than 12 h. On the other hand, the particle size distribution parameters of the milled W– 40 wt.% Ni powder mixtures increased at the early stage of milling, this is because Ni particles are more ductile than W. Thus, W particles were dispersed into the Ni ductile particles. This process delayed the fracturing of the W particles that was observed during the early stage of milling of pure W powders. The shift of the second peak to higher temperatures also indicates an increase in thermal stability of solid solutions formed at longer milling time. The first transformation peak temperatures of W–40% Ni powder mixtures were lower than those of pure W powders milled in the same conditions. The occurrence of lower transformation temperatures was attributed to wetting of W by Ni particles making it the crystal deformation of W easier. Nickel is more ductile and has a low melting temperature and thus enabling the formation of FCC Ni (W) solid solution upon high energy ball milling. Furthermore, it was noticed that the transformation temperatures of Ni(W) powders increased with milling time, which was ascribed to more atoms of W being dissolved in Ni lattices at longer milling time, hence separation
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of two exothermic peaks. Therefore, it infers from the variation of the transformation peak temperatures that the transformation temperature of pure Ni powder would decrease with milling time, whereas its mechanical alloying with refractory W would increase the thermal stability temperature of the milled powders. The formation of two peaks in the heating curve of W powder milled for 24 h and in that of W–40% Ni powders milled for 12 h indicates that two different phases were formed in each case out of the metastable products formed during ball milling. The appearance of the peaks in the curves of W–Ni mixtures milled for shorter times was ascribed to the faster kinetics of transformation in the mixtures than in pure W powders. Stronger peaks observed in the heating curves of powders milled for longer times (N24 h for W powders and N12 h in the case of W–Ni mixtures), thus showed that the homogeneity of the metastable products formed was improved and the dissolution between elemental powders was advanced at longer milling times. Annealing experiments were therefore conducted to identify the phase transformations taking place during heating and are discussed in Section 3.2. There was no decomposition peak of the stearic acid observed in the DTA curves. This could be attributed to the low content of the acid in the milled powders (3 wt.%) and to its loss during ball milling where an amount adhered onto the surfaces of the milling pot and balls. The enthalpy of decomposition of the stearic acid contained in the mixture of powders was therefore small comparatively to the latent heat of the samples analyzed. Hence the corresponding decomposition peaks were not observable. 3.2. Annealing transformations of milled powders The annealing temperatures were selected such that one was lower than the transformation peak temperature observed in the DTA
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curves (Fig. 1), the second temperature was 20 to 50 °C above the peak temperature and the third annealing temperature was 50 °C above the endothermic peak observed in the heating curve of milled W powder (Fig. 1a). Thus, milled powders were annealed at 550 °C, 730 °C and 1400 °C for 1 h in a tubular furnace in an argon flow. Scanning electron microscopy did not show any microstructural change of the W powder annealed at 550 °C, hence the peaks observed in the heating curves of W powders milled for 24 h (Fig. 1a) could be attributed to changes in crystal structure of metastable W as shown by XRD analysis (Fig. 3). The XRD patterns of W powder milled for 48 h and annealed at different temperatures (Fig. 3) indicated that HCP W was formed. Thus, metastable W induced by high energy mechanical milling of BCC W could transform into HCP W phase and the crystallites of the late grew upon annealing. Such transformation occurs through a shearing mechanism as in BCC to HCP similar to that of Fe under pressure or BCC to HCP in titanium below transition temperature of 883 °C [24]. Accordingly, during h i the shearing procedure, the (111)bcc act as habit plane, while 111 bcc act as shearing axis [25–27]. After shearing, some atoms were displaced from their original positions due to mechanical deformation; hence change in crystal orientation. The measured lattice parameters of the HCP phase were ah c p = 0.276 nm and c h c p = 0.481 nm. Using the equation V = a 2csin60° for HCP crystal, the a and c of HCP gives the density of 19.253 for HCP W. Some of W that did not transform to HCP due to lack of adequate deformation, transformed into FCC phase during annealing at 730 °C. As a result, FCC-based new eta-type phase with space group Fd-3 m (#227) and lattice parameter a = 1.109 nm has formed when a crystal network occurred between BCC space group #229 and metastable FCC W with space group #225. The possible mechanism for eta-type phase is such that the BCC lattice gets enveloped or gets bonded by two FCC lattices. As result, the lattice parameter of such superstructure is comprised of 2 × aFCC, plus 1 × aBCC, although thermal expansions must be taken into consideration. Thus the final lattice parameter must be approximately equal to 2 × 0.400 + 0.315 nm = 1.115 nm. The FCC and HCP phases were destroyed after annealing at 1400 °C. A BCC W single phase product
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was formed with slightly smaller lattice parameter as the XRD peaks had shifted by 0.3° toward high 2θ angles. Mechanical alloying led to the disordering of first the Ni crystal structure and its destruction before that of W as explained in [28]. It was also noticed that W peaks did not completely disappear after milling of W–Ni mixtures for 48 h (Fig. 4). However, these peaks vanished at low temperature annealing (~550 °C), a phenomenon attributed to thermally induced amorphization of mechanically alloyed powders [16,17]. The disappearance of W peak could also be ascribed to the thermal dissolution in nanoparticles of the Ni(W) solution. Two FCC phases were crystallized at high temperature annealing labeled (1) for the Ni-rich and (2) for the W-rich. Small peaks at about 550 °C were observed only in the DTA heating curves of the powders milled for 48 h (Fig. 1); no peaks were observed at low temperature in the curves of powders milled for shorter times. These peaks correspond to the disappearance from the X-ray diffraction patterns of the peaks of the residual W in the powders milled for 48 h as shown in Figs. 3 and 4. The small endothermic peaks formed at ~550 °C in the heating curves of W powders and W–Ni mixtures (Fig. 1) milled for 48 h could therefore attributed to a thermally induced amorphization of the residual W, transformation to a nanocrystalline W or its dissolution in W(Ni) solid solution. This low temperature transformation of mechanically milled powders was a precursor to the formation of various crystalline phases in both pure W and W–Ni powder mixtures. The sharp peak temperatures in DTA heating curves at ~ 730 °C are indicative of crystallization of milling products. The ICSD database [29] suggests the crystal structures of the products (1) and (2) formed above 730 °C be close to those of Ni10W and W3Ni2 with the lattice parameters 0.3577 nm and 1.1178 nm respectively. The lattice parameter of FCC Ni is 0.3519 nm and that of BCC W is 0.3194 nm before milling and annealing. Thus the crystal lattice of Ni expanded from 0.3519 nm to 0.3577 nm, which indicates dissolution of W with larger atomic diameter to form FCC solid solution and thus shifting of peaks toward lower diffraction angles. Thus the FCC solid solution formed a superstructure with Fd-3m (#227) space group by enveloping the untransformed W BCC lattice.
Fig. 3. XRD patterns of W powder, milled for 48 h and annealed at different temperatures for 1 h in an argon atmosphere. HCP W and FCC W were formed at intermediate annealing temperature and were destroyed at high temperature.
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Fig. 4. XRD patterns showing the disappearance of the residual W peak in W–Ni powder mixtures milled for 48 h annealed at 550 °C, and the formation of crystalline phases (1), (2) and (3) above 730 °C.
The discovery of a new phase “X” identified by the combination of SAD and CBD analysis of Ni–W thin films synthesized by magnetron sputtering was recently reported by Borgia et al. [30] in the center of
the phase diagram and β-tungsten in the tungsten-rich regime. The symmetry of the new X-phase was also identified as Fd-3m [30]. Similar to the case of pure W but in the case of W–Ni mixture, the FCC
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Fig. 5. BSE-SEM micrographs showing the microstructures of W–40 wt.% Ni powders milled for 48 h and annealed at 550 °C, 730 °C and 1400 °C with the formation and growth of W-rich phase (2) on the boundaries of the Ni-rich particles (1).
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Ni plays the role of binding BCC W lattices in forming W3Ni2 superstructure since presence of Ni enhanced the solid-state transformation in W, which would generate high compressive stress in the Ni matrix. The relaxation of the lattice strain in the Ni-rich matrix (1), therefore, contributes to the driving force for the formation of the W-rich product (2) on the boundaries of the particles of powders annealed at 730 °C (Fig. 5). The particles of the W-rich phase grew into faceted grains at high temperature (~1400 °C). Annealing below the peak temperatures (~550 °C) did not induce the formation of new phases in the powders (Fig. 5). The increase of the volume ratio of the W-rich phase (2) with temperature shows that more W-rich phase precipitated by decomposition of the metastable Ni(W) solid solution formed by milling. The XRD results (Fig. 4) show that upon annealing at 1400 °C, FCC Ni(W) solid solution dissociated, while the intensity of eta-phase (2) grew. Part of the metastable FCC W-rich phase transformed to HCP labeled (3). The intensity of the XRD peaks of the products (1) and (2) increased with the annealing temperature, between 730 °C and 1400 °C, due to the relaxation of lattice strains induced by ball milling and to the crystal growth. A significant grain growth was also reported by Fan et al. [31] in the mixtures of W, 7 wt.% Ni and 3 wt.% Fe milled for 20 h and annealed above 1100 °C. 3.3. Magnetic saturation of annealed powders The variation of the magnetic saturation (Ms) with milling time of pure W powder and that of W–Ni powders are illustrated in Fig. 6a. The corresponding saturation after annealing at 730 °C and 1400 °C are represented in Fig. 6b. The magnetic saturation of pure W powders increased with milling time in both milled and annealed conditions, whereas annealing at 730 °C and at 1400 °C yield the similar saturation values. The higher magnetic saturation of W was attributed to the occurrence of HCP W upon annealing as revealed by XRD analysis (Fig. 3). The increase in magnetic saturation after ball milling was discussed in [28], where it was shown that the BCC structure of W was deformed and
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transformed into HCP crystals. Thus, it infers from the XRD results and the magnetic saturation measurements that W powders became magnetic after mechanical milling and remained so after annealing. The relaxation of lattice strains at 730 °C and grain growth of W crystalline species (HCP, FCC and BCC) upon annealing did not suppress completely the magnetic character induced in the milled state powders. Therefore, crystal structure (packing density) after milling and annealing determined the Ms of pure W, whereas lattice strain and grain size played a minor role. However, annealing at high temperature (1400 °C) reverted the material to BCC structure, thus to the decrease in magnetic saturation value but not to zero as was observed in the initial W powder. The small magnetic character was therefore attributed to the deformation (residual compression) of the BCC structure of W after milling and annealing. On the other hand Ms of W–40 wt.% Ni powders decreased with milling time and annealing temperature. The lower magnetic saturation values of milled W–Ni powders annealed at 1400 °C than at 730 °C (Fig. 6b) were ascribed to ordering and grain growth of the magnetic matrix (1) at high temperature. Since the magnetic saturation in some conditions of milling and annealing is higher than in pure Ni as observed from unmilled mixture, the increased magnetic saturation could be attributed to the presence of anisotropic HCP W. Similarly, in pure W the increase in saturation is attributed to reverse transformation from metastable FCC to HCP, thus as temperature is increased more strain is relieved enabling it to revert back. This may be supported by Roebuck report on a linear decrease of the magnetic saturation of Co with the content of dissolved W during sintering [32]. Higher magnetic saturation of annealed biphase products than that of bulk powders had also been reported in NiZn ferrite/SiO2 composites [33]. 4. Conclusion The residual W contained in pure W and W–Ni mechanically milled powders underwent thermally induced transformation into HCP W or were dissolved in nanoparticles of Ni(W) solution during low temperature annealing before transformation in new crystalline phases at higher temperature.
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Subsequently, the nanocrystalline product formed by high energy ball milling of BCC W transformed into anisotropic HCP W and FCC W during high temperature annealing. The supersaturated Ni(W) solid solution formed during mechanical milling decomposed into two major products comprising of FCC Ni-rich and FCC W-rich products and formation of eta-type phase constituted of BCC lattice of W enveloped by two FCC lattices of Ni and W products. The structures and compositions of the two major phases were close to Ni10W and W3Ni2. The magnetic saturation of the milled W–Ni mixtures decreased with annealing temperature due to structural ordering and grain growth that follow the dissociation of supersaturated FCC Ni(W) and W(Ni) solutions into similar to Ni10W and W3Ni2. Acknowledgments This work was supported by the Diamond Research Laboratory at Element Six Ltd. Gratitude goes to Lancaster B., Maringa I., Mkhonto Z., Mukhali M., Mokoena E. and Thagisa J. References [1] Sagawa M, Fujimura S, Togawa M, Yamamoto H, Matsuura Y. New material for permanente magnets on a base of Nd and Fe (invited). J Appl Physiol 1984;55:2083–7. [2] Croat JJ, Herbst JF, Lee RW, Pinkerton FE. Pr–Fe and Nd–Fe-based materials: a new class of high performance permanent magnets (invited). J Appl Physiol 1984;55:2078–82. [3] Przybył A, Wysłocki JJ. The effect of annealing temperature on structure and magnetic properties of nanocomposite Nd10Fe84 − xWxB6 (0 b x b 33 at.%) magnets. J Mater Process Technol 2006;175:352–7. [4] Mondal BN, Basumallick A, Chattopadhyay PP. Effect of isothermal treatments on the magnetic behaviour of nanocrystalline Cu–Ni–Fe alloy prepared by mechanical alloying. J Magn Magn Mater 2007;309:290–4. [5] Rabanal ME, VáREZ A, Levenfeld B, Torralba JM. Magnetic properties of Mg-ferrite after milling process. J Mater Process Technol 2003;143–144:470–4. [6] Jartych E, Żurawicz JK, Oleszak D, Pękała M. Structure and magnetic properties of mechanosynthesized iron–tungsten alloys. J Magn Magn Mater 2000;218:247–55. [7] Besmel R, Shokrollahi H, Ghaffari M, Chitsazan B. Influence of milling time on the structural, microstructural and magnetic properties of mechanically alloyed Ni58Fe12Zr10Hf10B10 nanostructured/amorphous powders. J Magn Magn Mater 2011, doi:10.1016/j.jmmm.2011.05.025. [8] Bahgat M, Paek M-K, Pak J-J. Reduction investigation of WO3/NiO/Fe2O3 and synthesis of nanocrystalline ternary W–Ni–Fe alloy. J Alloys Compd 2009;472:314–8. [9] Osada M., Maeda K., Kawamoto Y., Shimizu Y., Nishimura T., Yamaharu S. US patent 2004; No. 6,773,247. [10] Metikos-Hukovic M, Grubac Z, Radic N, Tonejc A. Sputter deposited nanocrystalline Ni and Ni–W films as catalysts for hydrogen evolution. J Mol Catal A Chem 2006;249(1–2):172–80. [11] Genç A, Coşkun S, Öveçoğlu ML. Fabrication and properties of mechanically alloyed Ni activated sintered W matrix composites reinforced with T2O3 and TiB2 particles. Mater Charact 2010;21:740–8.
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Int. Journal of Refractory Metals and Hard Materials j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / I J R M H M
Thermal stability and magnetic saturation of annealed nickel–tungsten and tungsten milled powders Kasonde Maweja a,⁎, M.J. Phasha b, L.J. Choenyane a a b
Element Six Ltd/Diamond Research Laboratory, PO Box 561, Springs 1559, Gauteng, South Africa The Council for Scientific and Industrial Research, Materials Science and Manufacturing, Pretoria, South Africa
a r t i c l e
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Article history: Received 12 June 2011 Accepted 13 July 2011 Keywords: Nickel Tungsten Mechanical alloying Annealing Magnetic saturation
a b s t r a c t Crystal structures, microstructures and magnetic saturation of annealed pure W powder along with W–40 wt.% Ni powder mixtures processed by high-energy ball milling were investigated using XRD, DTA, SEM and saturation magnetization techniques. Thermally induced transformations occurred at low temperature annealing. Supersaturated metastable Ni(W) solid solution formed during mechanical milling decomposed during annealing treatment into FCC Ni-rich, FCC W-rich phases and an eta-type phase which was constituted of BCC lattice of W enveloped by two FCC lattices of Ni and W. The structures of the major annealing products were close to Ni10W and W3Ni2. The magnetic saturation of the milled W powder and W–Ni mixtures decreased with the increase in annealing temperature. Milling time was more influential on the magnetic properties of the annealed pure W powders. © 2011 Elsevier Ltd. All rights reserved.
1. Introduction Magnetic hard materials tend to find more and more applications in automotive industry (driving engines, clutches and frictionless bearings), in electronic (starters, sensors) and the like. A real breakthrough in the development of magnetic hard materials turned out to be the discovery, in 1984, of the Nd2Fe14B compound, which showed strong magnetocrystalline anisotropy characteristics [1,2]. The annealing temperature exerted significant impacts upon both structures and magnetic properties of the Nd–Fe–W–B magnets; in this case the effect of annealing temperature was attributed to the formation of Fe7W6 compounds at high annealing temperature [3]. Mondal et al. [4] showed that coarsening particles of mechanically alloyed Fe/FeNi powders impaired the magnetic properties, thus attempt of magnetic application of mechanically alloyed powders necessitates appropriate control over the precipitation/coarsening of the soft magnetic phases and temperature dependent magnetic response of the non magnetic matrix. Rabanal reported that the magnetic saturation of Mg-ferrite soft magnetic materials kept nearly constant, whereas the coercivity increased with milling time [5], whereas Jartych et al. [6] reported the decrease of magnetization versus milling time of mechanically alloyed Fe–W caused by the coexistence of the magnetic Fe(W) and the paramagnetic W(Fe). The studies of the magnetic properties showed that the coercivity decreased and saturation magnetization increased due to micro-stress elimination and reduction in surface dead layers and slight increase in
⁎ Corresponding author. Tel.: + 27 83 365 0952(cell). E-mail address: [email protected] (K. Maweja). 0263-4368/$ – see front matter © 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.ijrmhm.2011.07.005
ferromagnetic ordering upon annealing of Ni58Fe12Zr10Hf10B10 nanostructured mechanically alloyed powders [7]. Bahgat et al. [8] reached similar conclusion for W–Ni–Fe alloys prepared by mechanochemical reduction of mixtures corresponding oxide powders. Nickel–tungsten alloys have received attention as environmentally friendly substitute for hard chrome plating [9] and also because of their enhanced catalytic activity [10]. Activated sintering, that involves small additions of transition metals such as Ni, Co and Fe also enables major reductions in the sintering temperature of W [11–14]. XRD patterns of the sintered samples revealed the presence of (Ni,W) solid solution phase and a small peak of elemental Ni, whereas no intermetallic phases of Ni and W were formed during sintering [15]. On the other hand, thermally induced amorphization of the supersaturated solid solutions formed by mechanical alloying of elemental powders to form Zr52.5Cu17.9Ni14.6 Al10Ti5 and Zr70Pd30 products was observed by DSC and XRD analyses by El-Eskandarany et al. [16] and by Siegrist et al. [17] respectively. The amorphization process preceded the formation of crystalline phases upon annealing [17]. The degree of structural disorder induced by ball milling, the composition evolution upon annealing and the nanograin size distributions within the granular mixed ferromagnetic/anti-ferromagnetic (FM/AFM) powders are key factors for the magnetic properties in granular systems [18–20]. It has been reported that the magnetic saturation of Ni-based mechanically alloyed powders with Fe, Mo and Nb slightly increased at the early stage of milling before it decreased as a result of the electronic interactions between magnetic and non-magnetic elements and finally increased by the partial crystallization of the amorphous matrix [21]. By alloying the Ni–Fe alloys with ternary elements, such as Mo and changing the real structure from crystalline to the amorphous state, various kinds of
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magnetic, electrical and structural properties could be achieved [22,23]. Nickel is known to be a magnetic and tough material, whereas tungsten is hard but is not magnetic and has a high melting temperature (~ 3422 °C). Development of new tough magnetic structural materials for high temperature applications are therefore expected by alloying the two metals. The powder metallurgy route would present less technical challenges inherent to the high melting temperature of W for the manufacturing of this type of alloys. Hence, in the current study the effects of milling time and annealing temperature on the thermally induced transformations and magnetic saturation of W and W–Ni powders are reported.
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Vacuum gold sputter and were analyzed in the JEOL JSM-7500F Field Emission SEM equipped with EDX. Magnetic saturation of the annealed powders was determined using a commercial saturation induction measuring system (LDJ Electronics Inc.). The system was calibrated using pure cobalt and pure nickel references. The test piece of the milled powder was weighed and inserted into a non-magnetic holder. It was inserted into the high-intensity permanent magnet (~0.75 T), and then quickly withdrawn. The response of a search coil was displayed on a magnetic multimeter in terms of magnetic moment (saturation). Each measurement was repeated twice to check consistency of results. 3. Results and discussion
2. Material and experiments 3.1. Thermal analysis of milled powders Ball milling experiments of 100 g W powder at ~99.5% purity (1–2 μm) and 100 g powder mixtures of 60 wt.% W with 40 wt.% Ni (~99.5% purity, 2–3 μm) were conducted in a Retsch PM400 MA-type planetary high energy ball mill in an argon atmosphere at a rotation speed of 350 rpm. The grinding media consisted of 1 kg of 5 mm diameter 3Cr12 stainless steel balls. Stearic acid was added as 3 wt.% of total powder mass as process control agent to minimize cold welding. Milling times of 12 h, 24 h and 48 h were employed. The transformation temperatures of milled powders were determined by differential thermal analysis at a 10 °C/min heating rate in DTA-TGA equipment in an argon atmosphere. In order to characterize the thermally induced transformations, milled powders were annealed near the transformation temperatures for 1 h in a furnace filled with argon. The annealed products were then analyzed by XRD and SEM techniques. The annealed powders were cold mounted in epoxy resin and were mechanically polished using a series of SiC grit papers to cut through the cross-section of the powder particles. The metallographic polishing of samples was achieved using 1 μm diamond pastes. The samples were coated with gold using a Denton
The DTA heating curves of milled W powders and those of W–Ni mixtures are shown in Fig. 1(a) and (b) respectively. The peak temperatures of thermally induced transformation of the powders are indicated by the arrows. The non-milled powders of W and that of W–Ni mixtures were thermally stable throughout the temperature range used (25 °C to 1400 °C) as no transformation peak was observed in the respective heating curves. The curves of W powder (Fig. 1a) milled for 12 h had a small peak around 700 °C, whereas two small broad peaks were observed at ~680 °C and at ~730 °C in W powders milled for 24 h and a single sharp strong peak at ~720 °C was observed in the curves of powders milled for 48 h. All these peaks pointed toward exothermic character, which indicate that transformations, and not only strain relief, took place during the heating of milled W powders. On the other hand, the heating curves of W–40 wt.% Ni powder mixtures (Fig. 1b) milled for 12 h had two broad peaks at ~660 °C and 720 °C. The first exothermic peak shifted at ~ 670 °C for powders milled for 24 h and at around 680 °C for the W–Ni powders milled for
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48 h and its intensity was higher with increased milling time. In addition, the two exothermic peaks observed in W–40 wt.% Ni powders are further apart upon increased milling. First one positioned below 700 °C while the second lies above 800 °C as shown by the blue arrows in Fig. 1b. Due to inhomogeneity of ball milling, some fine and deformed W particles would transform earlier/faster than those that get exposed to severe deformations at a later stage of milling. The evolution of the particle sizes of pure W and W–40 wt.% Ni mixtures are shown in Fig. 2. The particle size distribution parameters of milled W powders presented minima at 12 h (Fig. 2a), which indicates that fracturing of the powder particles prevailed at the early milling stage. The d10, d50 and d90 of the powders were smaller than those of the starting powders and those of the products milled longer than 12 h. On the other hand, the particle size distribution parameters of the milled W– 40 wt.% Ni powder mixtures increased at the early stage of milling, this is because Ni particles are more ductile than W. Thus, W particles were dispersed into the Ni ductile particles. This process delayed the fracturing of the W particles that was observed during the early stage of milling of pure W powders. The shift of the second peak to higher temperatures also indicates an increase in thermal stability of solid solutions formed at longer milling time. The first transformation peak temperatures of W–40% Ni powder mixtures were lower than those of pure W powders milled in the same conditions. The occurrence of lower transformation temperatures was attributed to wetting of W by Ni particles making it the crystal deformation of W easier. Nickel is more ductile and has a low melting temperature and thus enabling the formation of FCC Ni (W) solid solution upon high energy ball milling. Furthermore, it was noticed that the transformation temperatures of Ni(W) powders increased with milling time, which was ascribed to more atoms of W being dissolved in Ni lattices at longer milling time, hence separation
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of two exothermic peaks. Therefore, it infers from the variation of the transformation peak temperatures that the transformation temperature of pure Ni powder would decrease with milling time, whereas its mechanical alloying with refractory W would increase the thermal stability temperature of the milled powders. The formation of two peaks in the heating curve of W powder milled for 24 h and in that of W–40% Ni powders milled for 12 h indicates that two different phases were formed in each case out of the metastable products formed during ball milling. The appearance of the peaks in the curves of W–Ni mixtures milled for shorter times was ascribed to the faster kinetics of transformation in the mixtures than in pure W powders. Stronger peaks observed in the heating curves of powders milled for longer times (N24 h for W powders and N12 h in the case of W–Ni mixtures), thus showed that the homogeneity of the metastable products formed was improved and the dissolution between elemental powders was advanced at longer milling times. Annealing experiments were therefore conducted to identify the phase transformations taking place during heating and are discussed in Section 3.2. There was no decomposition peak of the stearic acid observed in the DTA curves. This could be attributed to the low content of the acid in the milled powders (3 wt.%) and to its loss during ball milling where an amount adhered onto the surfaces of the milling pot and balls. The enthalpy of decomposition of the stearic acid contained in the mixture of powders was therefore small comparatively to the latent heat of the samples analyzed. Hence the corresponding decomposition peaks were not observable. 3.2. Annealing transformations of milled powders The annealing temperatures were selected such that one was lower than the transformation peak temperature observed in the DTA
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curves (Fig. 1), the second temperature was 20 to 50 °C above the peak temperature and the third annealing temperature was 50 °C above the endothermic peak observed in the heating curve of milled W powder (Fig. 1a). Thus, milled powders were annealed at 550 °C, 730 °C and 1400 °C for 1 h in a tubular furnace in an argon flow. Scanning electron microscopy did not show any microstructural change of the W powder annealed at 550 °C, hence the peaks observed in the heating curves of W powders milled for 24 h (Fig. 1a) could be attributed to changes in crystal structure of metastable W as shown by XRD analysis (Fig. 3). The XRD patterns of W powder milled for 48 h and annealed at different temperatures (Fig. 3) indicated that HCP W was formed. Thus, metastable W induced by high energy mechanical milling of BCC W could transform into HCP W phase and the crystallites of the late grew upon annealing. Such transformation occurs through a shearing mechanism as in BCC to HCP similar to that of Fe under pressure or BCC to HCP in titanium below transition temperature of 883 °C [24]. Accordingly, during h i the shearing procedure, the (111)bcc act as habit plane, while 111 bcc act as shearing axis [25–27]. After shearing, some atoms were displaced from their original positions due to mechanical deformation; hence change in crystal orientation. The measured lattice parameters of the HCP phase were ah c p = 0.276 nm and c h c p = 0.481 nm. Using the equation V = a 2csin60° for HCP crystal, the a and c of HCP gives the density of 19.253 for HCP W. Some of W that did not transform to HCP due to lack of adequate deformation, transformed into FCC phase during annealing at 730 °C. As a result, FCC-based new eta-type phase with space group Fd-3 m (#227) and lattice parameter a = 1.109 nm has formed when a crystal network occurred between BCC space group #229 and metastable FCC W with space group #225. The possible mechanism for eta-type phase is such that the BCC lattice gets enveloped or gets bonded by two FCC lattices. As result, the lattice parameter of such superstructure is comprised of 2 × aFCC, plus 1 × aBCC, although thermal expansions must be taken into consideration. Thus the final lattice parameter must be approximately equal to 2 × 0.400 + 0.315 nm = 1.115 nm. The FCC and HCP phases were destroyed after annealing at 1400 °C. A BCC W single phase product
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was formed with slightly smaller lattice parameter as the XRD peaks had shifted by 0.3° toward high 2θ angles. Mechanical alloying led to the disordering of first the Ni crystal structure and its destruction before that of W as explained in [28]. It was also noticed that W peaks did not completely disappear after milling of W–Ni mixtures for 48 h (Fig. 4). However, these peaks vanished at low temperature annealing (~550 °C), a phenomenon attributed to thermally induced amorphization of mechanically alloyed powders [16,17]. The disappearance of W peak could also be ascribed to the thermal dissolution in nanoparticles of the Ni(W) solution. Two FCC phases were crystallized at high temperature annealing labeled (1) for the Ni-rich and (2) for the W-rich. Small peaks at about 550 °C were observed only in the DTA heating curves of the powders milled for 48 h (Fig. 1); no peaks were observed at low temperature in the curves of powders milled for shorter times. These peaks correspond to the disappearance from the X-ray diffraction patterns of the peaks of the residual W in the powders milled for 48 h as shown in Figs. 3 and 4. The small endothermic peaks formed at ~550 °C in the heating curves of W powders and W–Ni mixtures (Fig. 1) milled for 48 h could therefore attributed to a thermally induced amorphization of the residual W, transformation to a nanocrystalline W or its dissolution in W(Ni) solid solution. This low temperature transformation of mechanically milled powders was a precursor to the formation of various crystalline phases in both pure W and W–Ni powder mixtures. The sharp peak temperatures in DTA heating curves at ~ 730 °C are indicative of crystallization of milling products. The ICSD database [29] suggests the crystal structures of the products (1) and (2) formed above 730 °C be close to those of Ni10W and W3Ni2 with the lattice parameters 0.3577 nm and 1.1178 nm respectively. The lattice parameter of FCC Ni is 0.3519 nm and that of BCC W is 0.3194 nm before milling and annealing. Thus the crystal lattice of Ni expanded from 0.3519 nm to 0.3577 nm, which indicates dissolution of W with larger atomic diameter to form FCC solid solution and thus shifting of peaks toward lower diffraction angles. Thus the FCC solid solution formed a superstructure with Fd-3m (#227) space group by enveloping the untransformed W BCC lattice.
Fig. 3. XRD patterns of W powder, milled for 48 h and annealed at different temperatures for 1 h in an argon atmosphere. HCP W and FCC W were formed at intermediate annealing temperature and were destroyed at high temperature.
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Fig. 4. XRD patterns showing the disappearance of the residual W peak in W–Ni powder mixtures milled for 48 h annealed at 550 °C, and the formation of crystalline phases (1), (2) and (3) above 730 °C.
The discovery of a new phase “X” identified by the combination of SAD and CBD analysis of Ni–W thin films synthesized by magnetron sputtering was recently reported by Borgia et al. [30] in the center of
the phase diagram and β-tungsten in the tungsten-rich regime. The symmetry of the new X-phase was also identified as Fd-3m [30]. Similar to the case of pure W but in the case of W–Ni mixture, the FCC
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Fig. 5. BSE-SEM micrographs showing the microstructures of W–40 wt.% Ni powders milled for 48 h and annealed at 550 °C, 730 °C and 1400 °C with the formation and growth of W-rich phase (2) on the boundaries of the Ni-rich particles (1).
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Ni plays the role of binding BCC W lattices in forming W3Ni2 superstructure since presence of Ni enhanced the solid-state transformation in W, which would generate high compressive stress in the Ni matrix. The relaxation of the lattice strain in the Ni-rich matrix (1), therefore, contributes to the driving force for the formation of the W-rich product (2) on the boundaries of the particles of powders annealed at 730 °C (Fig. 5). The particles of the W-rich phase grew into faceted grains at high temperature (~1400 °C). Annealing below the peak temperatures (~550 °C) did not induce the formation of new phases in the powders (Fig. 5). The increase of the volume ratio of the W-rich phase (2) with temperature shows that more W-rich phase precipitated by decomposition of the metastable Ni(W) solid solution formed by milling. The XRD results (Fig. 4) show that upon annealing at 1400 °C, FCC Ni(W) solid solution dissociated, while the intensity of eta-phase (2) grew. Part of the metastable FCC W-rich phase transformed to HCP labeled (3). The intensity of the XRD peaks of the products (1) and (2) increased with the annealing temperature, between 730 °C and 1400 °C, due to the relaxation of lattice strains induced by ball milling and to the crystal growth. A significant grain growth was also reported by Fan et al. [31] in the mixtures of W, 7 wt.% Ni and 3 wt.% Fe milled for 20 h and annealed above 1100 °C. 3.3. Magnetic saturation of annealed powders The variation of the magnetic saturation (Ms) with milling time of pure W powder and that of W–Ni powders are illustrated in Fig. 6a. The corresponding saturation after annealing at 730 °C and 1400 °C are represented in Fig. 6b. The magnetic saturation of pure W powders increased with milling time in both milled and annealed conditions, whereas annealing at 730 °C and at 1400 °C yield the similar saturation values. The higher magnetic saturation of W was attributed to the occurrence of HCP W upon annealing as revealed by XRD analysis (Fig. 3). The increase in magnetic saturation after ball milling was discussed in [28], where it was shown that the BCC structure of W was deformed and
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transformed into HCP crystals. Thus, it infers from the XRD results and the magnetic saturation measurements that W powders became magnetic after mechanical milling and remained so after annealing. The relaxation of lattice strains at 730 °C and grain growth of W crystalline species (HCP, FCC and BCC) upon annealing did not suppress completely the magnetic character induced in the milled state powders. Therefore, crystal structure (packing density) after milling and annealing determined the Ms of pure W, whereas lattice strain and grain size played a minor role. However, annealing at high temperature (1400 °C) reverted the material to BCC structure, thus to the decrease in magnetic saturation value but not to zero as was observed in the initial W powder. The small magnetic character was therefore attributed to the deformation (residual compression) of the BCC structure of W after milling and annealing. On the other hand Ms of W–40 wt.% Ni powders decreased with milling time and annealing temperature. The lower magnetic saturation values of milled W–Ni powders annealed at 1400 °C than at 730 °C (Fig. 6b) were ascribed to ordering and grain growth of the magnetic matrix (1) at high temperature. Since the magnetic saturation in some conditions of milling and annealing is higher than in pure Ni as observed from unmilled mixture, the increased magnetic saturation could be attributed to the presence of anisotropic HCP W. Similarly, in pure W the increase in saturation is attributed to reverse transformation from metastable FCC to HCP, thus as temperature is increased more strain is relieved enabling it to revert back. This may be supported by Roebuck report on a linear decrease of the magnetic saturation of Co with the content of dissolved W during sintering [32]. Higher magnetic saturation of annealed biphase products than that of bulk powders had also been reported in NiZn ferrite/SiO2 composites [33]. 4. Conclusion The residual W contained in pure W and W–Ni mechanically milled powders underwent thermally induced transformation into HCP W or were dissolved in nanoparticles of Ni(W) solution during low temperature annealing before transformation in new crystalline phases at higher temperature.
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Subsequently, the nanocrystalline product formed by high energy ball milling of BCC W transformed into anisotropic HCP W and FCC W during high temperature annealing. The supersaturated Ni(W) solid solution formed during mechanical milling decomposed into two major products comprising of FCC Ni-rich and FCC W-rich products and formation of eta-type phase constituted of BCC lattice of W enveloped by two FCC lattices of Ni and W products. The structures and compositions of the two major phases were close to Ni10W and W3Ni2. The magnetic saturation of the milled W–Ni mixtures decreased with annealing temperature due to structural ordering and grain growth that follow the dissociation of supersaturated FCC Ni(W) and W(Ni) solutions into similar to Ni10W and W3Ni2. Acknowledgments This work was supported by the Diamond Research Laboratory at Element Six Ltd. Gratitude goes to Lancaster B., Maringa I., Mkhonto Z., Mukhali M., Mokoena E. and Thagisa J. References [1] Sagawa M, Fujimura S, Togawa M, Yamamoto H, Matsuura Y. New material for permanente magnets on a base of Nd and Fe (invited). J Appl Physiol 1984;55:2083–7. [2] Croat JJ, Herbst JF, Lee RW, Pinkerton FE. Pr–Fe and Nd–Fe-based materials: a new class of high performance permanent magnets (invited). J Appl Physiol 1984;55:2078–82. [3] Przybył A, Wysłocki JJ. The effect of annealing temperature on structure and magnetic properties of nanocomposite Nd10Fe84 − xWxB6 (0 b x b 33 at.%) magnets. J Mater Process Technol 2006;175:352–7. [4] Mondal BN, Basumallick A, Chattopadhyay PP. Effect of isothermal treatments on the magnetic behaviour of nanocrystalline Cu–Ni–Fe alloy prepared by mechanical alloying. J Magn Magn Mater 2007;309:290–4. [5] Rabanal ME, VáREZ A, Levenfeld B, Torralba JM. Magnetic properties of Mg-ferrite after milling process. J Mater Process Technol 2003;143–144:470–4. [6] Jartych E, Żurawicz JK, Oleszak D, Pękała M. Structure and magnetic properties of mechanosynthesized iron–tungsten alloys. J Magn Magn Mater 2000;218:247–55. [7] Besmel R, Shokrollahi H, Ghaffari M, Chitsazan B. Influence of milling time on the structural, microstructural and magnetic properties of mechanically alloyed Ni58Fe12Zr10Hf10B10 nanostructured/amorphous powders. J Magn Magn Mater 2011, doi:10.1016/j.jmmm.2011.05.025. [8] Bahgat M, Paek M-K, Pak J-J. Reduction investigation of WO3/NiO/Fe2O3 and synthesis of nanocrystalline ternary W–Ni–Fe alloy. J Alloys Compd 2009;472:314–8. [9] Osada M., Maeda K., Kawamoto Y., Shimizu Y., Nishimura T., Yamaharu S. US patent 2004; No. 6,773,247. [10] Metikos-Hukovic M, Grubac Z, Radic N, Tonejc A. Sputter deposited nanocrystalline Ni and Ni–W films as catalysts for hydrogen evolution. J Mol Catal A Chem 2006;249(1–2):172–80. [11] Genç A, Coşkun S, Öveçoğlu ML. Fabrication and properties of mechanically alloyed Ni activated sintered W matrix composites reinforced with T2O3 and TiB2 particles. Mater Charact 2010;21:740–8.
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