SEM and Elemental Mapping study of mechanically alloyed binary W-Ni

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S. Behera, S. S. Khuntia, A. Patra*

*Nanomaterials Research Group, Department of Metallurgical and Materials Engineering, National Institute of Technology, Rourkela, Rourkela-769008, Odisha, India

Anshuman Patra is presently working as an Assistant Professor in Metallurgical and Materials Engineering Department, National Institute of Technology, Rourkela, Orissa, India since 2014. He was the recipient of university medal and Indranil award from MGMI Kolkata in B.E. He has six years of industrial experience in various fields such as production and quality division of Ductile Iron pipe Industry, Research and Development Department of Aluminum Rolling Industry. He is the author of the book on aluminium alloy processing and coauthor of a book on process modeling for steel industry. His scientific interests include nanomaterials, spark plasma sintering, advanced high temperature materials and ODS alloys.
Tungsten (W) and Nickel (Ni) with nominal composition of 90W10Ni (alloy A), 905W5Ni (alloy B), (in weight %) has been mechanically alloyed for 20 h followed by compaction at 500 MPa for 5 mins with three varying thickness from each composition. The compacted pellets are subjected to conventional pressureless sintering at 1500 °C for two hours under continuous Ar flow. The phase formation, microstructure mechanical and high temperature oxidation behavior at 1000 °C for 10 hours of the alloys are studied. Formation of Ni-W intermetallic phase at W matrix interface offers significant pinning action in alloys with minimum thickness and hinders the grain boundary migration. Sintered alloy A with higher Ni content and specimen with minimum thickness exhibit maximum densification (86.79%), hardness (604.14 HV), wear resistance and superior oxidation resistance. The present study will provide a roadmap to adjust the sample dimension and nickel content for high temperature application of W based alloys.
This work has been supported by TEQIP, NIT Rourkela.
Corresponding author
Prof. Anshuman Patra, Metallurgical and Materials Engineering Department, National Institute of Technology-Rourkela
Rourkela-769008, Orissa.
Office +91 661-2462574
W is the most desirable refractory material for several high temperature structural applications owing to its high melting point, superior density, hardness and strength1. However the formation of volatile oxides at elevated temperature results in severe degradation of the component and radioactive emission in fusion reactor2.
An increase in environment temperature beyond 1000 °C causes serious diminution of mass (150 kg/h) due to excessive formation of volatile species (WO3)3. The phenomenon restricts the application of W either < 500 °C or in vacuum atmosphere2. Ni addition is quite significant with respect to low cost and enhancement of densification of W based alloys which in turn has immense importance to improve the mechanical and high temperature oxidation resistance.
According to literature evidence contiguity, % segment and mean free path between the W grains stimulates the mechanical behavior of W alloys4. Consolidation of W alloys by conventional electrical resistance route in hydrogen environment urges higher temperature (≥ 2500 °C) which is not encouraging due to the problem of excessive grain growth and cost factors5-7.
Synthesis of W-Ni alloys by conventional melting-casting is not attractive due to low solubility of Ni in W and wide variation in melting point of W (3410 °C) and Ni (1455 °C)8. Fabrication of nanostructured W alloys with the addition of suitable activator such as Ni leads to substantial reduction of sintering temperature and boosting of sintering kinetics. Sintering at slightly higher temperature as compared to melting point of Ni offers additional driving force as capillary pressure of formed liquid phase and therefore enhances atomic bonding. 
Pellet thickness after compaction may influence the homogeneity of the component. Higher thickness of compacted pellet requires more applied pressure and dwell time for uniform pressure transmission throughout the pellet. Increased pressure enhances the plastic deformation of powders followed by increase in green density of the pellet9, 10. However, higher pressure application (1 GPa) during compaction of nano-W powder is problematic in view of increased strength, surface area and brittleness11. Therefore in our present investigation 500 MPa pressure is employed for compaction. 
Compacted pellet thickness possesses significant importance in order to achieve enhance densification and compositional homogenization. Increase in pressure enhances the density of the pellet before sintering by increased particle deformation12. Sample with higher thickness required enhanced compaction stress to achieve equivalent densification as compared lower thickness sample13. Therefore present investigation deals with phase, microstructure evolution, mechanical and high temperature behavior of W-Ni alloys with varied thickness.
Commercial W and Ni powders are mechanically alloyed for 20 h in a high energy planetary ball mill (Fritsch, P5) with 300 rpm, toluene as process control agent (PCA) and 10 mm diameter grinding steel balls. The two compositions have been selected as 90%W, 10% Ni (alloy A) and 95%W, 5% Ni (alloy B). Powders are removed at an interval of 1 h, 10 h, 20 h for characterization.
Scanning electron microscopy (SEM) [JEOL, JSM-6084LV] with accelerating voltage: 20 KV, secondary electron imaging (SEI) mode has been used for studying the morphology of milled powder and Field emission scanning electron microscopy (FESEM) (FEI Novanano SEM 450) with accelerating voltage: 15 KV and backscattered imaging (CBS) has been employed for studying the sintered alloy and oxide samples.
Environmental scanning electron microscopy (ESEM) (FEI Quanta FEG 250) has been used for analysis of wear track morphology and microstructure of oxidized samples. Elemental mapping (EM) and energy dispersive spectroscopy (EDS) has been employed for compositional analysis of the sintered alloys, oxidized samples.
The 20 h milled powders are placed in a 10 mm diameter die-punch (tool steel made) and 500 MPa pressure has been applied for fabrication of cylindrical pellets.
The thickness of the compacted pellets is provided in table 1. The compacted pellets after compaction are presented in Fig. 1.

Table 1.
 Top, Thickness of compacted pellet
Figure 1. Above, Thickness variation of compacted pellets of (a) alloy A, (b) alloy B.
The compacted pellets are subjected to sintering at 1500 °C for two hours holding time in a tubular furnace under constant Argon flow to prevent sample oxidation. The samples are placed on a refractory plate and pushed inside the furnace for sintering and cooled inside the furnace.
The green density (pellet density before sintering) and sintered density has been measured by Archimedes principle14
The hardness of the samples are measured by micro-hardness tester (Leco, LM 248AT) with 100 gf load and 10 s dwell time. 
To comprehend the abrasion characteristics wear study has been conducted by ball on plate wear tester (Ducom, TR-208-M1) for 10 minutes at 25 rpm speed and 20 N load. Evenness of both top and bottom surface has been checked before wear study.
The high temperature oxidation behavior has been studied in a raising hearth furnace at 1000 °C for 10 hours. The W based alloys are used in kinetic energy penetrator and rocket nozzle and therefore evaluation of oxidation resistance is utmost importance to assesses the high temperature efficiency.
The microstructure of the mechanically alloyed powders (alloy A and alloy B) at different milling time is investigated by SEM (Fig. 2(a-d) and Fig. 3(a-d)). The powder particles show uneven size and morphology at 0 h. Continued deformation during milling leads to reduction in the particle size and flattening of particles. The size reduction between solvent (W) and solute (Ni) (relatively lesser extent) minimizes the distance for diffusion and therefore Ni atoms jump in W lattice for formation of an alloy. Maximum particle refinement takes place between 1-10 h of milling when the defect generation is intense. The particle flattening (flake shape formation) is more predominant at 10 and 20 h of milling. Along with flattening, nearly spherical and irregular particle morphology are also observed at 20 h of milling. Several instances of finer particle coalescence are also evident owing to requirement of higher strength of fracturing of finer particles15. Mechanical alloying improves the packing factor of particles owing to morphological modifications with reduction of interconnection between particles16. The particle size distribution in Fig. 4 (a, b) shows that alloy A exhibits an increasingly finer particle size and improved particle bimodality (combination of finer and coarser particles) as compared to alloy B which influences the compaction and further sintering behavior.

Figure 2
. Left, SEM of alloy A at (a) 0 h, (b) 1 h, (c) 10 h, (d) 20 h of milling

Figure 3
. Right, SEM of alloy B at (a) 0 h, (b) 1 h, (c) 10 h, (d) 20 h of milling.

Figure 4
. Variation of particle size of 20 h milled (a) alloy A, (b) alloy B.
Fig. 5(a-c) and Fig. 6(a-c) shows the microstructure of sintered samples A1, A2, A3 and B1, B2, B3 respectively. The bright, gray, and combined black-gray phases directs to the W matrix, Ni rich matrix and Ni-W intermetallics phases. Few porosities are also observed and indicated by arrows. The phases such as Ni rich matrix in sample B1 is identified by point EDS analysis as displayed in Fig. 7 (a, b). Other phases in sample A1 is identified similarly. Elemental maps of sample A1 and B1 indicate the presence of Ni (red dots) inside W matrix due to diffusion of Ni atoms in W lattice in course of mechanical alloying (Fig. 8, Fig. 9). Higher Ni rich area of sample A1 (of alloy A) against sample B1 (of alloy B) is accredited to higher Ni weight% in alloy A. Few spherical and few elongated shape grain morphologies are evident in sample A1 designated as the incidence of non-uniform liquid phase sintering process. The liquid phase sintering process also contributes in closing of open porosity17. Dissolution of W grains in Ni grains and later crystallization of intermetallic compound along with flowing of liquid Ni to the porosities leads to improve the density of the samples12.
The intermetallics are mostly located at the W/W matrix and W/Ni matrix interface. The intermetallics inhibit the grain boundary mobility and may yield in fine grain structure. No oxide formation is evident in the microstructure owing to the inert environment during sintering. The grain size distribution of the sintered sample (Fig. 10) reveals that the maximum grain refinement occurs in sample A1 as compared to other samples. 

Figure 5
. Microstructure of sintered sample (a) A1, (b) A2, (c) A3 of 20 h milled alloy A.
Figure 6. Microstructure of sintered sample (a) B1, (b) B2, (c) B3 of 20 h milled alloy B.
Figure 7. (a) SEM image, (b) corresponding EDS pattern on pointed area of sintered sample B1.
Figure 8. (a) SEM image, corresponding elemental maps of (b) W, (c) Ni of sintered sample A1.
Figure 9. (a) SEM image, corresponding elemental maps of (b) W, (c) Ni of sintered sample B1.
Figure 10. Grain size distribution in sintered sample (a) A1, B1, (b) A2, B2, (c) A3, B3.
Porosity/Density has a significant impact on the mechanical properties of a material. The % relative green density [(green density/theoretical density)*100] and % sintered density enhances with decreasing thickness. 
As discussed earlier, alloy A shows improved distribution of finer and coarser particles which results in increase compaction behavior, closing of pores followed by further improvement of particle bonding and increase density during sintering phase.
Maximum % relative green density and % relative sintered density of 56.46%, 86.79% respectively has been recorded for sample A1. Lower Ni content in alloy B leads to formation of reduced liquid Ni matrix and closing of porosity against alloy A. In addition, higher thickness pellets urges higher compaction pressure and dwell time to achieve equivalent density as compared to thin pellets.
Therefore, the reduced density in thick pellets may be accredited to non-homogeneous pressure distribution with the same load and time.
The hardness value is influenced by the W grain size, compositional distribution, and presence of porosity. Reduced porosity, finer grain size and higher presence of W may give rise to higher hardness18.
Formation of solid solution also impacts the hardness of an alloy. Atomic radius of Ni (0.149 nm) is substantially less than W (0.193 nm) and develops an extensive strain field which interact with the dislocation stress field and contributes in hardening of the alloy19.
Precipitation of Ni-W intermetallic compound at the W matrix interface also minimizes the grain boundary mobility during sintering and refines the microstructure. The hardness variation shows the similar trend as that of density variation for sample A1, A2, A3 (Fig. 11). Maximum mean hardness of 604.14 HV (Vickers Hardness) has been achieved for sample A1 owing to improved density and fine grain size. 
Figure 11. Hardness of sintered sample A1, A2, A3 of alloy A.
The hardness and abrasion resistance of hard materials such as W is effective for its application in cutting tool and rocket nozzle, kinetic energy penetrator applications.
The wear morphology of sample A1, A3 and sample B1, B3 is presented in Fig. 12(a-d). It is observed from the wear study that the wear loss increases with increasing thickness. Minimum wear width of 258.43 µm has been achieved in sample A1as compared to A3 (271.72 µm), B1 (281 µm) and B3 (307.89 µm).
Figure 12. Left, Microstructure of wear track for sintered sample (a) A1, (b) A3, (c) B1, (d) B3.
The material flow at the edges is non-uniform as well as scratches on wear tracks are more in high thickness samples. Several presences of wear particles in sample B3 represents scratching followed by cracking of surface.
Significant scratching reflects less plastic deformation ability of the sample. Presence of porosity and low hardness also enhance the wear loss. The porosity may aggravate the scratching and further cracking phenomenon which leads to higher wear volume. Therefore, alloy A exhibit better wear resistance owing to high hardness and enhanced densification than alloy B.
The oxidation kinetics of alloy A (sample A1, A2, A3) and alloy B (sample B1, B2, B3) is presented in Fig. 13(a, b). It is apparent that initially the oxidation kinetics follows linear trend, however alloy B shows steep increase in weight change per unit area as compared to shallow increase in alloy A.
Figure 13. Below, Oxidation kinetics of sintered sample (a) A1, A2, A3 and (b) B1, B2, B3 at 1000 °C for 10 h. 
The constant rise in weight change is ascribed to the formation of porous volatile WO3 without any continuous barrier along the interface18, 20. During oxidation, grain boundary diffusion is dominant unless covered by protective scale and therefore cracking along the interface will further enhance the intensity of oxidation. The ratio of oxygen and tungsten points to the vacancy concentration in oxidized samples.
Higher ratio suggests more vacancies and easy pathways of incoming oxygen and outgoing volatile oxide products18. The difference between molar volume of WO3 and NiWO4 is considerably less than WO3 and WO220. Therefore, less W and more Ni content in alloy A facilitates the reduction in WO3 content and increase in dense NiWO4 concentration which may be responsible for enhancement of oxidation resistance in alloy A against alloy B. As discussed previously, vacancies such as pores causes augmentation in oxidation rate, minimum porosity (maximum density) in sample A1 exhibit superior oxidation resistance against rest of the samples.
Further, microstructure of oxide samples displays porous microstructure and cracking of oxide layer with increase in sample thickness (from A1 to A3) (Fig. 14(a-c)). Oxide microstructure of sample A1 (Fig. 14(a))  reveals more compactness of oxide scale in sample A1 as compared to other samples (Fig.14 (b-d)). EDS results (Fig. 15(a, b)) reveals lower W and higher Ni content in sample A1 against sample B1 which indicates lower porous WO3 and increased dense NiWO4 oxide formation in sample A1.
The final oxide sample images (Fig. 16(a-f)) submits that blistering and spallation increases with increase in sample thickness and in alloy B as compared to alloy A as a whole. No serious cracking is evident in oxidized sample A1 against other samples. It is manifested from present investigation that sample A1 possesses superior physical (density), mechanical (hardness, wear) and elevated temperature oxidation resistance against rest of the samples.

Figure 14
. Above, Microstructure of 10 h oxidized sample of (a) A1, (b) A2, (c) A3 and (d) B1 at 1000 °C.
Figure 15. EDS spectra of 10 h oxidized sample (a) A1, (b) B1 at 1000 °C.
Figure 16. Oxidized sample (10 h) image of (a) A1, (b) A2, (c) A3, (d) B1, (e) B2, (f) B3.
W-Ni alloys with varied composition and thickness have been fabricated by mechanical alloying followed by compaction and sintering. The study shows enriched diffusion of Ni in W during mechanical alloying. Microstructure of sintered alloys shows the incidence of liquid phase formation and finer grain size is achieved in 90W10Ni alloy (alloy A). Superior density, hardness and wear resistance is recorded in 90W10Ni alloy with minimum thickness. High temperature oxidation study displays that oxidation resistance enhances with decreasing thickness and superior hindrance against oxidation is observed in 90W10Ni alloy.
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