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Microstructure and mechanical properties of tungsten heavy alloys
Materials Science and Engineering A 527 (2010) 7841–7847
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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure and mechanical properties of tungsten heavy alloys Jiten Das a,∗ , G. Appa Rao a , S.K. Pabi b a b
Defence Metallurgical Research Laboratory, Hyderabad, India Indian Institute of Technology, Kharagpur, India
a r t i c l e
i n f o
Article history: Received 22 June 2010 Received in revised form 20 August 2010 Accepted 24 August 2010
Keywords: Transgranular Intergranular Fractographs Tungsten heavy alloy
a b s t r a c t Two tungsten (W)–nickel (Ni)–copper (Cu) alloys (WNCs) and one W–Ni–iron (Fe) alloy (WNF) were prepared by liquid phase sintering at 1783 K and 1733 K, respectively. The average W-grain size in the sintered WNCs (60–70 m) was coarser than that in the WNF alloy (30 m) possibly due to the higher sintering temperature (1783 K) necessary for the former alloys. The volume of the matrix phase in the WNF (25–30 vol.%) was higher than that in WNCs (10–15 vol.%). The tensile properties and hardness of WNF specimens at room temperature were significantly superior to those of WNC specimens apparently due to the finer W-grain size, lesser contiguity and porosity in the former. WNF specimens, in contrast to WNCs, failed under tension by W-grain cleavage fracture, possibly due to relatively stronger matrix phase and W/matrix bonding. At very low strain rate (0.0001/s) the tensile curve of WNF was wavy in nature, but these were absent at higher strain rates (0.001 to 1/s). The tensile strength and elongation of WNF alloy remarkably deteriorated at higher temperatures (773 and 973 K), and the fracture changed to matrix failure mode apparently due to weakening of the matrix phase. © 2010 Elsevier B.V. All rights reserved.
1. Introduction W heavy alloys (WHAs) are conventionally produced by liquid phase sintering method. In this process, W powder is mixed with relatively low melting elemental powders such as Ni, Fe, Cu, Co, etc., compacted by either hydraulic press or cold isostatic press (CIP), and sintered in a furnace with a continuous flow of hydrogen. During sintering, the lower melting elements melt and form the matrix that bonds the unmelted W particles together. Some amount of W also gets dissolved in the matrix and gets re-precipitated on the primary W particles, which renders the W-grains rounded and larger in size [1]. Thus, the sintered microstructure of WHAs essentially consists of rounded W grains cemented by a ductile and relatively low melting matrix phase [1]. WHAs with W content 90–95% possess good combination of tensile properties (tensile strength ∼1000 MPa and elongation ∼30%), as well as, high density of 17–18 g/ml [2]. Because of their high density, these alloys find wide applications as centre of gravity (CG) adjuster, radiation shields, kinetic energy penetrators (KEP), etc. [3–5]. In general, the properties of WHAs depend on their microstructural features. Properties of WHAs can also be tailored by controlling the size and the volume fraction of the W grains [6], and also by controlling the strength of the W/matrix interface [7,8]. Zhou et al. [9] reported that relatively higher strength,
∗ Corresponding author. Tel.: +91 9866867563; fax: +91 4024344535. E-mail address: [email protected] (J. Das). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.08.071
toughness, and elongation values (883 MPa, 29 J and 10%) were observed for 95W–5(Ni/Fe) alloy, when the interfacial bonding strength between binder phase (matrix) and W grain was higher. Stronger interfacial bonding strength renders ductile fracture of the matrix and cleavage fracture of the W grains [8,9]. The strength of interface between the W grain and the matrix is largely governed by the cleanliness of the W powder, composition of the matrix and the W-grain size. Proper post-sintering heat treatment, however, avoids segregation of impurities at the tungsten–matrix interface and allows the impurities to remain homogeneously distributed throughout the matrix [2]. Therefore, the post-sintering heat treatment causes great improvement in mechanical properties of WHAs [9]. Alloy chemistry also affects the mechanical properties of WHAs. For example, in sintered conditions the W–Ni–Cu base heavy alloys have much inferior tensile properties and hardness at room temperature as compared to those of W–Ni–Fe base heavy alloys [10]. The properties of WHAs are found to be dependent on the amount of W content [11,12], and on Ni/Fe [18,11] or Ni/Cu ratio [10]. Song et al. [11] reported that the heavy alloys with Ni/Fe ratio of 4:1 shows best tensile properties and impact strength. However, German and Bourguignon [2] reported that the Ni/Fe ratio of 7:3 was optimal. Bose and Kapoor [12] showed that lowering the W content in the alloys up to 90 wt% caused an improvement in the ultimate tensile strength as well as elongation value of the sintered WHAs. They also observed that 92.5%W–(Ni–Fe–Co) alloy deformed up to 95% showed a tensile strength of 1720 MPa with concomitant elongation of 16%. Katavik and Nikacevic [13] explained that
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Table 1 Nominal and analyzed composition of WHAs prepared in this investigation. Alloy
WNC1 WNC2 WNF
alloy was also designed and produced in this study. Although, W–Ni–Cu base heavy alloys possess inferior tensile properties and hardness at room temperature as compared to those of W–Ni–Fe base heavy alloys [10], proper explanation in this regard was not available in the literature. This work was undertaken to understand the compositional effect also. Limited high temperature tensile properties of 92.5W–5Ni–2.5 Fe alloy under very restricted test condition (strain rate 0.000083/s, partly under high vacuum), was available [16] in the literature. The high temperature tensile properties of WHAs – tested in air specifically at 500 ◦ C and above – are not readily available in the literature. However, this understanding is important because, when a WHA projectile hits a barrier, its temperature rapidly increases as it continues to deform. Therefore, the WNF alloy – produced in this study – was subjected to tensile test both at room temperature as well as at elevated temperatures (500 and 700 ◦ C) in air at different strain rates in order to generate data for tensile properties of this alloy at elevated temperature.
Composition, wt%
Nominal Analyzed Nominal Analyzed Nominal Analyzed
W
Ni
Cu
95.00 94.89 96.00 96.10 91.00 91.04
3.5 3.42 3.00 2.80 7.00 6.86
1.5 1.69 1.00 1.10
Fe
Co
1.50 1.45
0.50 0.60
the mechanical properties of W–Ni–Co with nickel to cobalt ratios ranging from 2 to 9 are far superior to that of W–Ni–Fe alloys. Bose and Kapoor [12] also showed that WHAs doped with Mo and Re exhibit improved yield strength. However, according to the study of Liu et al. [14], higher Mo addition to WHAs leads to brittleness due to the formation of a precipitate phase. Although, Re is very effective strengthener for WHAs [12], its use is very restricted because of its very high cost. Based on the above literature and earlier work carried out by the authors [8,15], suitable alloys were designed to obtain correlation between chemistry, microstructure and mechanical properties of WHAs. Three alloys namely WNC1 (W—95.0 wt%, Ni—3.5 wt% and Cu—1.5 wt%), WNC2 (W—96.0 wt%, Ni—3.0 wt% and Cu—1.0 wt%), and WNF (W—91.0 wt%, Ni—7.0 wt%, Fe—1.5 wt%, and Co—0.50 wt%) alloy were designed and produced by conventional powder metallurgical route. Here, Co is added to dissolve more W in to the matrix and thereby, act as matrix strengthener in WNF alloy [13]. WNF alloy was expected to have best combination of tensile properties and WNC2 alloy was expected to have worst combination of tensile properties. Furthermore, according to German and Bourguignon [2], major change in elongation values were observed in the W–Ni–Fe based heavy alloys with W content ranging between 95 and 96 wt%. In order to verify whether the same effect is observed in W–Ni–Cu based heavy alloys or not, WNC1
2. Experimental Two W–Ni–Cu base heavy alloys, namely WNC1 (95W–3.5wt%Ni–1.5wt%Cu alloy) and WNC2 (96W–3wt%Ni– 1wt%Cu alloy), and one W–Ni–Fe base heavy alloy WNF (91W–7wt%Ni–1.5w%Fe–0.5wt%Co alloy) were prepared by liquid phase sintering technique. The particle size distribution of W powder used for making the elemental blends was measured by Malvern particle size analyzer. W powder was admixed with relatively low melting powders such as Ni, Fe, Cu, Co. in the required proportions and cold isostatically pressed (Make: National Forge, Belgium: maximum pressure 520 MPa) at 250 MPa pressure. Liquid phase sintering temperature is mainly dependent on the W content in WHAs. The higher the W content in the WHAs, the higher is the required sintering temperature. Optimum liquid phase sintering temperatures for different WHAs were determined earlier by trial and error method [15]. The WNF and WNCs alloys were liq-
Table 2 Density and tensile properties and hardness of WHAs prepared in this study vis-à-vis the properties of WHAs obtained from the literature. Alloy
W-grain size (m)
Densification %
Tensile strength (MPa)
El. (%)
Contiguity
WNC1
60
98.4
660 ± 10 (strain rate 1/s)
3
0.60
WNC2
70
98.4
660 ± 12 (strain rate 1/s)
3
0.70
WNF
30
99.4
1000 ± 20 (strain rate 1/s)
20
0.33
93W–4.9Ni–2.1Fe (sintered at 1530 ◦ C/70 min) Vac. annealing at 1200 ◦ C for 2 h [17] 90W–7Ni–3Fe (sintered at 800 ◦ C/1 h, 1500 ◦ C/30 min) Heat treatment 1100 ◦ C/1 h/WQ [18] 90W–7Ni–3Fe (attritor milled, CIPed, 230 MPa, sintered at 1480 ◦ C/0.5 h) [19] 90W–7Ni–3Fe (attritor milled, CIPed, 230 MPa, sintered at 1480 ◦ C/0.5 h) [20] 90W–7Ni–3Fe [2]
30
960
23
17
∼950
∼30
0.25
∼19
99.6
∼970
∼25
∼0.28
∼19
99.6
∼1000
∼30
∼0.32
∼1020
∼36
∼0.28
95W–5(Ni/Fe) as sintered condition [9]
883
10
95W–5(Ni/Fe = 7/3) as sintered condition [1]
991
14
∼0.5
916
11
∼0.65
95W–3.5Ni–1.5Fe (sintered 1500 ◦ C for 30 min and vacuum annealed at 1100 ◦ C for 8 h) [20]
∼99
(1) Bulk hardness (2) Micro-hardness of W grain (3) Micro-hardness of matrix phase in GPa (1) 2.76 ± 0.2 (2) 4.50 ± 0.3 (3) 2.50 ± 0.3 (1) 2.86 ± 0.2 (2) 4.48 ± 0.2 (3) 2.70 ± 0.3 (1) 3.20 ± 0.1 (2) 4.33 ± 0.3 (3) 3.10 ± 0.2 Bulk hardness 28–29 HRC
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uid phase sintered at 1733 and 1783 K, respectively for 1.5 h in H2 atmosphere in a continuous pusher type furnace (Make: FHD Furnace Limited, England). The heavy alloys are not generally sintered for more than 1.5 h in order to avoid pore coarsening by Ostwald ripening [2]. After that the alloys were heat treated at 1373 K for 5 h in argon (Ar) atmosphere and quenched in water. The Ar atmosphere helps in removal of absorbed H2. The water quenching is done to avoid segregation of impurities at the phase boundaries and to allow the impurities to remain homogeneously distributed [2]. The tensile properties of WNF alloy before and after post-sintering heat treatment were evaluated and tensile fracture surface before and after heat treatment was recorded. The heat treated WNF alloy was also subjected to swaging with an intermediate annealing at 500 ◦ C for 1 h in order to observe the effect of the swaging operation on tensile properties of WNF alloy. The densities of the alloys were measured using Archimedes principle. The optical microscopy was carried out by Leitz Optical Microscope. The X-ray mappings of elements like W, Ni, Cu and Fe were done using an electron probe microanalyzer (EPMA) (Make: Sx-100, Cameca, France). The matrix composition of all the alloys was also analyzed by EPMA. Hardnesses of the W grains and the binding phase (matrix) were measured separately by micro-hardness tester (Make: MATSUZAWA, Japan). The bulk hardness of the alloys was measured by Wolpert Vicker Hardness Tester using 30 kg load. The tensile specimens were prepared from these alloys as per the ASTM standard E 8 (2003) and they were tested using a Masanto Tensometer (maximum capacity: 2000 kg). The fractured surfaces of the tensile tested specimens were observed using LEO scanning electron microscope (SEM). Room temperature (298 K) and elevated temperatures (773 K, and 973 K) tensile tests of WNF alloy were carried out in air by means of an INSTRON universal testing machine (maximum capacity 10000 kg) and the corresponding tensile stress–strain behaviour was recorded. The fractured surfaces of these specimens were examined by SEM. Energy dispersive spectroscopy (EDS) of the fracture surface of these specimen was also performed. 3. Results and discussion The nominal and analyzed (wet chemical method) composition of WNC1, WNC2 and WNF alloys are shown in Table 1. These alloys were prepared from the W powder with 10 m average particle size. The densities and mechanical properties of the alloys after sintering and heat treatment are recorded in Table 2. The tensile properties and fracture surfaces of WNF alloy before and after post-sintering heat treatment and after swaging are recorded in Table 3 and Fig. 1 respectively in order to study the effect postsintering heat treatment operation and swaging operation on the tensile properties and tensile fracture behaviour of the alloy. The WNF alloy showed significant improvement of tensile properties after the post-sintering heat treatment operation (Table 3). The tensile fracture mode changes from predominantly matrix or interface fracture (before heat treatment) to transgranular tungsten grain cleavage fracture after the post-sintering heat treatment operation (Fig. 1). During sintering preferential segregation of impurities Table 3 Tensile properties of WNF alloy before and after post-sintering heat treatment operation and after swaging operation. Properties
Before post-sintering heat treatment
After post-sintering heat treatment
After swaging up to 15% reduction in area (with an intermediate annealing at 500 ◦ C for 1 h)
Tensile strength Elongation
717 MPa 5%
1000 MPa 20%
1167 18%
Fig. 1. Tensile fracture surfaces of (a) sintered, (b) sintered and heat treated and (c) sintered, heat treated and swaged WNF alloy respectively.
takes place at the matrix and W-grain interface and H2 absorption takes place [2]. The segregation of these impurities weakens the interface [2] while H2 causes embrittlement [20]. During tensile testing of the sintered alloys, interface failure/matrix failure takes place which is associated with the lower tensile properties [8]. However, post-sintering operation prevent embrittlement [20] by de-absorption of residual H2 [2] and by allowing the impurities to remain homogeneously distributed throughout the matrix. This causes matrix and interface stronger and failure takes place predominantly by W-grain cleavage failure which is associated with higher tensile properties [8]. Therefore, all the alloys prepared in
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Fig. 2. Optical microstructures of (a) WNC1, (b) WNC2 and (c) WNF alloy after postsintering heat treatment.
this study were heat treated before evaluating their microstructure and mechanical properties. The heat treated samples were then swaged up to 15% reduction in area with an intermediate annealing at 500 ◦ C for 1 h. It was observed that the tensile properties improved significantly (T.S −1167 MPa, elongation 18%) (Table 3) after swaging treatment. This may be due to significant reduction of W-grain size of the WNF alloy after this treatment. After swaging treatment, the WNF alloy failed during tensile test predominantly by W-grain cleavage fracture (Fig. 1) which is associated with higher tensile properties [8]. Optical microstructures of the sintered and heat treated WHAs are displayed in Fig. 2. The volume of the matrix phase in the sintered and heat treated WNF (25–30 vol.%) was higher than that in WNC1 (10–15 vol.%) and WNC2 (10–15 vol.%) alloy. Average W-grain sizes of WNF, WNC1 and WNC2 alloy, were observed to be about 30 m, 60 m and 70 m respectively. Because of the slightly higher sintering temperature (1783 K) employed for the WNC alloy as compared to that for WNF alloy, the W-grain size was observed to be slightly higher in the former alloys. More shape accommodation and larger size W grains, observed in WNCs, is also due to high amount of W content in these alloys. [2]. Because, the smaller the separation between the W grains (as in case of higher W containing alloys), the higher is the grain growth observed during sintering [2]. Table 2 shows that the percentage densification observed in case of the WNC1 and WNC2 alloys was almost identical (98.4%), whereas that for WNF was still higher (99.4%). The higher amount of intergranular (matrix) phase
Fig. 3. Tensile fractographs of the sintered and heat treated (a) WNC1, (b) WNC2 and (c) WNF alloy respectively. The tensile test was carried out at room temperature and at a strain rate of 1/s.
present in the WNF, as compared to that in WNC1 and WNC2 alloys, apparently contributed to this better densification of WNF alloy. On the other hand WNCs possess insufficient liquid for densification [2]. The micro-hardnesses of the matrix phase (binding phase) and W grains as well as the bulk hardness and the tensile properties of the alloys after the post-sintering heat treatment operation are summarized in Table 2. The bulk hardness of the WNC alloys is found to be lower than that of WNF alloy. It is also observed that the micro-hardness of the W-grains are identical in all three alloys, while the micro-hardness of the matrix phase of WNF alloy (3.1 GPa) is slightly higher than that of WNC alloys (2.5–2.7 GPa). The difference in hardness, however, is much less pronounced than the difference in tensile strength and elongation values (Table 2) of WNF and WNCs. The tensile properties of WNF alloy are far superior to those of the WNC alloys. This might be due to the fact that W-grain size, and contiguity of WNF alloy are smaller as compared to those of WNC alloys [2] as shown in Table 2. The tensile fracture surfaces of both the WNC alloys and WNF specimen are shown in Fig. 3. While predominantly matrix or interface failure is observed in the case of WNCs alloys, the mode of failure is predominantly Wgrain cleavage fracture in WNF alloy at room temperature. Since the micro-hardness of the W grains in WNF and WNCs is similar, and the micro-hardness of the matrix phase in WNF is only slightly higher than that in WNC (Table 2), the vastly higher tensile
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Fig. 4. W, Ni, Fe and Cu elemental maps for (a) WNF, (b) WNC1 and (c) WNC2 alloy respectively.
strength of the WNF as compared to that of WNCs and the transgranular cleavage fracture in WNF possibly originated from finer W-grain size and better W/matrix phase cohesion and lesser contiguity (0.33) in WNF alloy as compared to those of WNC alloys. The other reported values [2,17–19] of microstructural parameters such as contiguity, W-grain size and mechanical properties are listed in Table 2 for comparison. The difference in contiguity between this study and literature is due to clustering of tungsten grains [1]. Where as the strength values are comparable, the difference in elongation values between this study and the literature is due to relatively higher contiguity and larger W-grain size of the WNF alloy produced in this study. WNC1 (95W–3.5Ni–1.5Cu) produced in this study showed inferior tensile properties (T.S. −660 MPa, El. −3%) as compared to 95W–5(Ni/Fe) alloy properties available in the literature [1,9,20]. It is interesting to note from the EPMA analysis that the matrix of the WNF alloy is richer in W (∼35 wt%) as compared to that of WNC alloys (∼20–23.4 wt%) (Table 4). W, Ni, Fe and Cu elemental maps for WNF, WNC1 and WNC2 alloy respectively are shown in Fig. 4. Some amount of porosity (marked by arrow in Fig. 3a and b)
was observed in the tensile fracture surfaces of WNC alloys but not in that of WNF alloy, which conformed well with the higher level of densification achieved in case of WNF alloy (Table 2). Lower tensile properties of WNC alloys might be partly due to the presence of these porosities. As WNF alloys show better room temperature tensile properties further room temperature tensile testing at different strain rates (0.1, 0.01, 0.001, 0.0001/s) were carried out and the results of these tests are displayed in Fig. 5. Analysis of these data show (Table 5) that in general at room temperature the yield stress and tensile stress of WNF increases and the % elongation diminishes with the increase in strain rate, which is expected. However, the mode of
Table 4 EPMA analysis (wt%) of the matrix phase of WHAs. Alloy
W
Ni
Cu
WNC1 WNC2 WNF
23.4 ± 1.9 20.2 ± 2.6 34.9 ± 6.5
58.8 ± 2.4 57.7 ± 1.4 49.9 ± 7.1
18.4 ± 2.5 23.7 ± 3.8
Fe
9.6 ± 1.4
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Table 5 Tensile properties of WNF at room temperature and at elevated temperatures vis-à-vis the room temperature and elevated temperature tensile properties of WHAs obtained from the literature. Alloy
Tensile test temperature(◦ C)
Strain rate (s−1 )
Yield stress (MPa)
WNF (tested in air)
25 25 25 25 500 700 25 500 700
0.1 0.01 0.001 0.0001 0.00003 0.000067 0.000083 0.000083 0.000083
922 803 716 630 368 370 ∼800 ∼450 ∼400
92.5W–5Ni–2.5Fe alloy [16] (tested partly under high vacuum)
Fig. 5. Tensile stress–strain curves of room temperature and elevated temperature tested samples which were prepared from sintered and heat treated WNF alloy.
fracture remained trans granular at all strain rates. The room temperature stress–strain curve at slow strain rate of 0.0001/s is wavy in nature; but this is absent at higher strain rates (Fig. 5). Hence, wavy nature in the tensile curve is not due to the presence of any porosity in the sintered and heat treated WNF alloy. In fact, no such microporosity was revealed by the micrograph in Fig. 3c. The wavy nature of the tensile curve indicates inhomogeneous deformation. The dislocation locking–unlocking phenomenon at low strain rate might be responsible for rendering the wavy tensile curve [21]. Room temperature tensile properties are comparable with available data in the literature [16]. The tensile properties of WNF alloy was also evaluated at elevated temperatures (773 K and 973 K), and the results are shown
Engineering stress at maximum load (MPa) 1054 1045 994 966 548 427 ∼1050 ∼700 ∼600
Total elongation (%) 25.4 26.8 37.5 35.1 4.90 3.00 ∼20 ∼15 ∼16
in Table 5 and Fig. 5. The results, when compared with room temperature test data (Table 5), show that both tensile strength and yield strength decreases by more than 40% at elevated temperatures, but a catastrophic decrease in the % elongation of the sample by more than 85% occurred upon rising the temperature. However, the catastrophic decrease in % elongation is not observed [16] at elevated temperatures (773 K and 973 K) even at a very slow strain rates (0.000083/s), when the WHA alloy specimens are tensile tested partly under high vacuum. Since 973 K is far lower than the recrystallization temperature of W (>1650 K [22]), the deterioration of the tensile properties at 973 or 773 K must be due to softening and embrittlement of the matrix phase by oxidation during the test conducted in air. Oxidation (Fig. 6b) on the fracture surface has occurred due to the fact that tensile test was carried out at a very slow strain rate in air. The fractograph of the high temperature tensile tested specimen is shown in Fig. 6a. Here the specimen has failed by predominantly matrix fracture during tensile testing, whereas at room temperature WNF failed predominantly by W-grain cleavage fracture (Fig. 3c). 4. Conclusions The tensile properties and hardness of WNF alloy at room temperature are superior due to the presence of finer W-grain size, lesser contiguity and porosity in its microstructure. WNF alloy fails under tension by W-grain cleavage fracture due to relatively stronger matrix phase and W/matrix bonding. The tensile properties and hardness of WNC alloys at room temperature are inferior due to the presence of coarse W-grain size, higher contiguity and porosity in their microstructure. WNC alloys fail under tension by matrix/interface fracture due to relatively weaker matrix phase and W/matrix bonding.
Fig. 6. (a) Fractograph of broken surface of WNF alloy which was tensile tested at 773 K and at a strain rate of 0.00003/s and (b) EDS graph obtained from the fracture surface.
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When WNF alloy is tested in air for prolonged time at elevated temperatures (773 and 973 K), its tensile strength and elongation remarkably deteriorates and the fracture changes to matrix failure mode due to weakening of the matrix phase. Acknowledgements Authors convey their sincere gratitude to Dr. G Malakondaiah, Distinguished Scientist and Director, DMRL, for giving permission to publish this work. The financial support provided by DRDO is thankfully acknowledged. They would like to thank the staff members of Powder Metallurgy Group, DMRL for their help in experimentation. Thanks are also due to Shri Bijoy Sarma, Division Head, Powder Metallurgy Division, for necessary support and useful discussion. One of the authors (JD) is indebted to Mr. A Saibabu, Sc. ‘F’, (retired) for his helpful suggestions and guidance to carry out this investigation. Authors gratefully acknowledge Shri A Sambasiva Rao for his support during EPMA work. References [1] R. Gero, L. Borukhin, I. Pikus, Mater. Sci. Eng. A 302 (2001) 162–167. [2] R.M. German, L.L. Bourguignon, Powder Metall. Def. Technol. 6 (1984) 117–131. [3] F.V. Lenel, Powder Metallurgy Principles and Application, MPIF, Princeton, 1980. [4] R.M. German, L.L. Bourguignon, B.H. Rabin, Powder Metall. 5 (1992) 3–13. [5] R. Gero, D. Chaiat, in: I. Minkoff (Ed.), High Density Tungsten Alloys, Mater. Eng. Conf., Israel, 1981, pp. 46–50.
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[6] S.H. Islam, X.H. Qu, F. Akhtar, P.Z. Feng, X.B. He, Mater. Sci. Forum 534–36 (2007) 561–564. [7] W. Liu, Y. Ma, B. Huang, Bull. Mater. Sci. 31 (2008) 1–6. [8] J. Das, U.R. Kiran, A. Chakraborty, N.E. Prasad, Int. J. Refract. Met. Hard Mater. 27 (2009) 577–583. [9] W. Zhou, X. Gao, Y.-G. Zhou, F.-H. Luo, Mater. Sci. Eng. Powder Metall. 15 (2010) 141–144. [10] Information Available at Refractory Materials & Ceramics Branch, Advanced Technology & Materials Co. Ltd., No. 76, Xueyuan, Nanlu, Haidian, Beijing 100081, China, Email: [email protected] (downloaded from www.refractorymetal.net on June 2, 2008). [11] Song, H.-S. Kim, E.-P. Lee, Seong, The Effect of Ni/Fe Ratio on the Mechanical Properties for Tungsten Heavy Alloy, From Advanced Materials, Manufacturing and Testing Analysis Centre, 1998, Downloaded from http://ammtiac.alionscience.com/ammt/iacdocs.do?M005374 on 29May-2008. [12] A. Bose, D. Kapoor, in: A. Bose, R.J. Dowding (Eds.), Tungsten, Refractory Metals and Alloys 4: Processing, Properties and Applications, 1998. [13] B. Katavik, M. Nikacevic, Properties of the Cold Swaged and Strain Aged P/M 91W–6Ni–3Co Heavy Alloy, Second International Conference on: Deformation Processing and Structure of Materials, Belgrade, Serbia and Montenegro, 2005, pp. 135–140. [14] W. Liu, Y. Ma, B. Huang, Bull Mater. Sci. 31 (2008) 1–6. [15] J. Das, Un-published work, 2007. [16] E. Pink, S. Kumar, Mater. Sci. Eng. A 234–236 (1997) 102–105. [17] Y. Yu, L. Hu, E. Wang, Mater. Sci. Eng. A 435–436 (2006) 620–624. [18] B.H. Rabin, A. Bose, R.M. German, Int. J. Powder Metall. 25 (1989) 21–26. [19] S. Eroglu, T. Baykara, J. Mater. Process. Technol. 103 (2000) 288–292. [20] S.H. Islam, F. Akhtar, S.J. Askari, M. Tufail, X. Qu, Int. J. Refract. Met. Hard Mater. 25 (2007) 380–385. [21] F.R.N. Nebarro, J.P. Hirth, Dislocation in Solids, Elsevier, Oxford, 2004. [22] S.N. Mathaudhu, A.J. deRosset, K.T. Hartwig, L.J. Kecskes, Mater. Sci. Eng. A 503 (2009) 28–31.
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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure and mechanical properties of tungsten heavy alloys Jiten Das a,∗ , G. Appa Rao a , S.K. Pabi b a b
Defence Metallurgical Research Laboratory, Hyderabad, India Indian Institute of Technology, Kharagpur, India
a r t i c l e
i n f o
Article history: Received 22 June 2010 Received in revised form 20 August 2010 Accepted 24 August 2010
Keywords: Transgranular Intergranular Fractographs Tungsten heavy alloy
a b s t r a c t Two tungsten (W)–nickel (Ni)–copper (Cu) alloys (WNCs) and one W–Ni–iron (Fe) alloy (WNF) were prepared by liquid phase sintering at 1783 K and 1733 K, respectively. The average W-grain size in the sintered WNCs (60–70 m) was coarser than that in the WNF alloy (30 m) possibly due to the higher sintering temperature (1783 K) necessary for the former alloys. The volume of the matrix phase in the WNF (25–30 vol.%) was higher than that in WNCs (10–15 vol.%). The tensile properties and hardness of WNF specimens at room temperature were significantly superior to those of WNC specimens apparently due to the finer W-grain size, lesser contiguity and porosity in the former. WNF specimens, in contrast to WNCs, failed under tension by W-grain cleavage fracture, possibly due to relatively stronger matrix phase and W/matrix bonding. At very low strain rate (0.0001/s) the tensile curve of WNF was wavy in nature, but these were absent at higher strain rates (0.001 to 1/s). The tensile strength and elongation of WNF alloy remarkably deteriorated at higher temperatures (773 and 973 K), and the fracture changed to matrix failure mode apparently due to weakening of the matrix phase. © 2010 Elsevier B.V. All rights reserved.
1. Introduction W heavy alloys (WHAs) are conventionally produced by liquid phase sintering method. In this process, W powder is mixed with relatively low melting elemental powders such as Ni, Fe, Cu, Co, etc., compacted by either hydraulic press or cold isostatic press (CIP), and sintered in a furnace with a continuous flow of hydrogen. During sintering, the lower melting elements melt and form the matrix that bonds the unmelted W particles together. Some amount of W also gets dissolved in the matrix and gets re-precipitated on the primary W particles, which renders the W-grains rounded and larger in size [1]. Thus, the sintered microstructure of WHAs essentially consists of rounded W grains cemented by a ductile and relatively low melting matrix phase [1]. WHAs with W content 90–95% possess good combination of tensile properties (tensile strength ∼1000 MPa and elongation ∼30%), as well as, high density of 17–18 g/ml [2]. Because of their high density, these alloys find wide applications as centre of gravity (CG) adjuster, radiation shields, kinetic energy penetrators (KEP), etc. [3–5]. In general, the properties of WHAs depend on their microstructural features. Properties of WHAs can also be tailored by controlling the size and the volume fraction of the W grains [6], and also by controlling the strength of the W/matrix interface [7,8]. Zhou et al. [9] reported that relatively higher strength,
∗ Corresponding author. Tel.: +91 9866867563; fax: +91 4024344535. E-mail address: [email protected] (J. Das). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.08.071
toughness, and elongation values (883 MPa, 29 J and 10%) were observed for 95W–5(Ni/Fe) alloy, when the interfacial bonding strength between binder phase (matrix) and W grain was higher. Stronger interfacial bonding strength renders ductile fracture of the matrix and cleavage fracture of the W grains [8,9]. The strength of interface between the W grain and the matrix is largely governed by the cleanliness of the W powder, composition of the matrix and the W-grain size. Proper post-sintering heat treatment, however, avoids segregation of impurities at the tungsten–matrix interface and allows the impurities to remain homogeneously distributed throughout the matrix [2]. Therefore, the post-sintering heat treatment causes great improvement in mechanical properties of WHAs [9]. Alloy chemistry also affects the mechanical properties of WHAs. For example, in sintered conditions the W–Ni–Cu base heavy alloys have much inferior tensile properties and hardness at room temperature as compared to those of W–Ni–Fe base heavy alloys [10]. The properties of WHAs are found to be dependent on the amount of W content [11,12], and on Ni/Fe [18,11] or Ni/Cu ratio [10]. Song et al. [11] reported that the heavy alloys with Ni/Fe ratio of 4:1 shows best tensile properties and impact strength. However, German and Bourguignon [2] reported that the Ni/Fe ratio of 7:3 was optimal. Bose and Kapoor [12] showed that lowering the W content in the alloys up to 90 wt% caused an improvement in the ultimate tensile strength as well as elongation value of the sintered WHAs. They also observed that 92.5%W–(Ni–Fe–Co) alloy deformed up to 95% showed a tensile strength of 1720 MPa with concomitant elongation of 16%. Katavik and Nikacevic [13] explained that
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Table 1 Nominal and analyzed composition of WHAs prepared in this investigation. Alloy
WNC1 WNC2 WNF
alloy was also designed and produced in this study. Although, W–Ni–Cu base heavy alloys possess inferior tensile properties and hardness at room temperature as compared to those of W–Ni–Fe base heavy alloys [10], proper explanation in this regard was not available in the literature. This work was undertaken to understand the compositional effect also. Limited high temperature tensile properties of 92.5W–5Ni–2.5 Fe alloy under very restricted test condition (strain rate 0.000083/s, partly under high vacuum), was available [16] in the literature. The high temperature tensile properties of WHAs – tested in air specifically at 500 ◦ C and above – are not readily available in the literature. However, this understanding is important because, when a WHA projectile hits a barrier, its temperature rapidly increases as it continues to deform. Therefore, the WNF alloy – produced in this study – was subjected to tensile test both at room temperature as well as at elevated temperatures (500 and 700 ◦ C) in air at different strain rates in order to generate data for tensile properties of this alloy at elevated temperature.
Composition, wt%
Nominal Analyzed Nominal Analyzed Nominal Analyzed
W
Ni
Cu
95.00 94.89 96.00 96.10 91.00 91.04
3.5 3.42 3.00 2.80 7.00 6.86
1.5 1.69 1.00 1.10
Fe
Co
1.50 1.45
0.50 0.60
the mechanical properties of W–Ni–Co with nickel to cobalt ratios ranging from 2 to 9 are far superior to that of W–Ni–Fe alloys. Bose and Kapoor [12] also showed that WHAs doped with Mo and Re exhibit improved yield strength. However, according to the study of Liu et al. [14], higher Mo addition to WHAs leads to brittleness due to the formation of a precipitate phase. Although, Re is very effective strengthener for WHAs [12], its use is very restricted because of its very high cost. Based on the above literature and earlier work carried out by the authors [8,15], suitable alloys were designed to obtain correlation between chemistry, microstructure and mechanical properties of WHAs. Three alloys namely WNC1 (W—95.0 wt%, Ni—3.5 wt% and Cu—1.5 wt%), WNC2 (W—96.0 wt%, Ni—3.0 wt% and Cu—1.0 wt%), and WNF (W—91.0 wt%, Ni—7.0 wt%, Fe—1.5 wt%, and Co—0.50 wt%) alloy were designed and produced by conventional powder metallurgical route. Here, Co is added to dissolve more W in to the matrix and thereby, act as matrix strengthener in WNF alloy [13]. WNF alloy was expected to have best combination of tensile properties and WNC2 alloy was expected to have worst combination of tensile properties. Furthermore, according to German and Bourguignon [2], major change in elongation values were observed in the W–Ni–Fe based heavy alloys with W content ranging between 95 and 96 wt%. In order to verify whether the same effect is observed in W–Ni–Cu based heavy alloys or not, WNC1
2. Experimental Two W–Ni–Cu base heavy alloys, namely WNC1 (95W–3.5wt%Ni–1.5wt%Cu alloy) and WNC2 (96W–3wt%Ni– 1wt%Cu alloy), and one W–Ni–Fe base heavy alloy WNF (91W–7wt%Ni–1.5w%Fe–0.5wt%Co alloy) were prepared by liquid phase sintering technique. The particle size distribution of W powder used for making the elemental blends was measured by Malvern particle size analyzer. W powder was admixed with relatively low melting powders such as Ni, Fe, Cu, Co. in the required proportions and cold isostatically pressed (Make: National Forge, Belgium: maximum pressure 520 MPa) at 250 MPa pressure. Liquid phase sintering temperature is mainly dependent on the W content in WHAs. The higher the W content in the WHAs, the higher is the required sintering temperature. Optimum liquid phase sintering temperatures for different WHAs were determined earlier by trial and error method [15]. The WNF and WNCs alloys were liq-
Table 2 Density and tensile properties and hardness of WHAs prepared in this study vis-à-vis the properties of WHAs obtained from the literature. Alloy
W-grain size (m)
Densification %
Tensile strength (MPa)
El. (%)
Contiguity
WNC1
60
98.4
660 ± 10 (strain rate 1/s)
3
0.60
WNC2
70
98.4
660 ± 12 (strain rate 1/s)
3
0.70
WNF
30
99.4
1000 ± 20 (strain rate 1/s)
20
0.33
93W–4.9Ni–2.1Fe (sintered at 1530 ◦ C/70 min) Vac. annealing at 1200 ◦ C for 2 h [17] 90W–7Ni–3Fe (sintered at 800 ◦ C/1 h, 1500 ◦ C/30 min) Heat treatment 1100 ◦ C/1 h/WQ [18] 90W–7Ni–3Fe (attritor milled, CIPed, 230 MPa, sintered at 1480 ◦ C/0.5 h) [19] 90W–7Ni–3Fe (attritor milled, CIPed, 230 MPa, sintered at 1480 ◦ C/0.5 h) [20] 90W–7Ni–3Fe [2]
30
960
23
17
∼950
∼30
0.25
∼19
99.6
∼970
∼25
∼0.28
∼19
99.6
∼1000
∼30
∼0.32
∼1020
∼36
∼0.28
95W–5(Ni/Fe) as sintered condition [9]
883
10
95W–5(Ni/Fe = 7/3) as sintered condition [1]
991
14
∼0.5
916
11
∼0.65
95W–3.5Ni–1.5Fe (sintered 1500 ◦ C for 30 min and vacuum annealed at 1100 ◦ C for 8 h) [20]
∼99
(1) Bulk hardness (2) Micro-hardness of W grain (3) Micro-hardness of matrix phase in GPa (1) 2.76 ± 0.2 (2) 4.50 ± 0.3 (3) 2.50 ± 0.3 (1) 2.86 ± 0.2 (2) 4.48 ± 0.2 (3) 2.70 ± 0.3 (1) 3.20 ± 0.1 (2) 4.33 ± 0.3 (3) 3.10 ± 0.2 Bulk hardness 28–29 HRC
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uid phase sintered at 1733 and 1783 K, respectively for 1.5 h in H2 atmosphere in a continuous pusher type furnace (Make: FHD Furnace Limited, England). The heavy alloys are not generally sintered for more than 1.5 h in order to avoid pore coarsening by Ostwald ripening [2]. After that the alloys were heat treated at 1373 K for 5 h in argon (Ar) atmosphere and quenched in water. The Ar atmosphere helps in removal of absorbed H2. The water quenching is done to avoid segregation of impurities at the phase boundaries and to allow the impurities to remain homogeneously distributed [2]. The tensile properties of WNF alloy before and after post-sintering heat treatment were evaluated and tensile fracture surface before and after heat treatment was recorded. The heat treated WNF alloy was also subjected to swaging with an intermediate annealing at 500 ◦ C for 1 h in order to observe the effect of the swaging operation on tensile properties of WNF alloy. The densities of the alloys were measured using Archimedes principle. The optical microscopy was carried out by Leitz Optical Microscope. The X-ray mappings of elements like W, Ni, Cu and Fe were done using an electron probe microanalyzer (EPMA) (Make: Sx-100, Cameca, France). The matrix composition of all the alloys was also analyzed by EPMA. Hardnesses of the W grains and the binding phase (matrix) were measured separately by micro-hardness tester (Make: MATSUZAWA, Japan). The bulk hardness of the alloys was measured by Wolpert Vicker Hardness Tester using 30 kg load. The tensile specimens were prepared from these alloys as per the ASTM standard E 8 (2003) and they were tested using a Masanto Tensometer (maximum capacity: 2000 kg). The fractured surfaces of the tensile tested specimens were observed using LEO scanning electron microscope (SEM). Room temperature (298 K) and elevated temperatures (773 K, and 973 K) tensile tests of WNF alloy were carried out in air by means of an INSTRON universal testing machine (maximum capacity 10000 kg) and the corresponding tensile stress–strain behaviour was recorded. The fractured surfaces of these specimens were examined by SEM. Energy dispersive spectroscopy (EDS) of the fracture surface of these specimen was also performed. 3. Results and discussion The nominal and analyzed (wet chemical method) composition of WNC1, WNC2 and WNF alloys are shown in Table 1. These alloys were prepared from the W powder with 10 m average particle size. The densities and mechanical properties of the alloys after sintering and heat treatment are recorded in Table 2. The tensile properties and fracture surfaces of WNF alloy before and after post-sintering heat treatment and after swaging are recorded in Table 3 and Fig. 1 respectively in order to study the effect postsintering heat treatment operation and swaging operation on the tensile properties and tensile fracture behaviour of the alloy. The WNF alloy showed significant improvement of tensile properties after the post-sintering heat treatment operation (Table 3). The tensile fracture mode changes from predominantly matrix or interface fracture (before heat treatment) to transgranular tungsten grain cleavage fracture after the post-sintering heat treatment operation (Fig. 1). During sintering preferential segregation of impurities Table 3 Tensile properties of WNF alloy before and after post-sintering heat treatment operation and after swaging operation. Properties
Before post-sintering heat treatment
After post-sintering heat treatment
After swaging up to 15% reduction in area (with an intermediate annealing at 500 ◦ C for 1 h)
Tensile strength Elongation
717 MPa 5%
1000 MPa 20%
1167 18%
Fig. 1. Tensile fracture surfaces of (a) sintered, (b) sintered and heat treated and (c) sintered, heat treated and swaged WNF alloy respectively.
takes place at the matrix and W-grain interface and H2 absorption takes place [2]. The segregation of these impurities weakens the interface [2] while H2 causes embrittlement [20]. During tensile testing of the sintered alloys, interface failure/matrix failure takes place which is associated with the lower tensile properties [8]. However, post-sintering operation prevent embrittlement [20] by de-absorption of residual H2 [2] and by allowing the impurities to remain homogeneously distributed throughout the matrix. This causes matrix and interface stronger and failure takes place predominantly by W-grain cleavage failure which is associated with higher tensile properties [8]. Therefore, all the alloys prepared in
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Fig. 2. Optical microstructures of (a) WNC1, (b) WNC2 and (c) WNF alloy after postsintering heat treatment.
this study were heat treated before evaluating their microstructure and mechanical properties. The heat treated samples were then swaged up to 15% reduction in area with an intermediate annealing at 500 ◦ C for 1 h. It was observed that the tensile properties improved significantly (T.S −1167 MPa, elongation 18%) (Table 3) after swaging treatment. This may be due to significant reduction of W-grain size of the WNF alloy after this treatment. After swaging treatment, the WNF alloy failed during tensile test predominantly by W-grain cleavage fracture (Fig. 1) which is associated with higher tensile properties [8]. Optical microstructures of the sintered and heat treated WHAs are displayed in Fig. 2. The volume of the matrix phase in the sintered and heat treated WNF (25–30 vol.%) was higher than that in WNC1 (10–15 vol.%) and WNC2 (10–15 vol.%) alloy. Average W-grain sizes of WNF, WNC1 and WNC2 alloy, were observed to be about 30 m, 60 m and 70 m respectively. Because of the slightly higher sintering temperature (1783 K) employed for the WNC alloy as compared to that for WNF alloy, the W-grain size was observed to be slightly higher in the former alloys. More shape accommodation and larger size W grains, observed in WNCs, is also due to high amount of W content in these alloys. [2]. Because, the smaller the separation between the W grains (as in case of higher W containing alloys), the higher is the grain growth observed during sintering [2]. Table 2 shows that the percentage densification observed in case of the WNC1 and WNC2 alloys was almost identical (98.4%), whereas that for WNF was still higher (99.4%). The higher amount of intergranular (matrix) phase
Fig. 3. Tensile fractographs of the sintered and heat treated (a) WNC1, (b) WNC2 and (c) WNF alloy respectively. The tensile test was carried out at room temperature and at a strain rate of 1/s.
present in the WNF, as compared to that in WNC1 and WNC2 alloys, apparently contributed to this better densification of WNF alloy. On the other hand WNCs possess insufficient liquid for densification [2]. The micro-hardnesses of the matrix phase (binding phase) and W grains as well as the bulk hardness and the tensile properties of the alloys after the post-sintering heat treatment operation are summarized in Table 2. The bulk hardness of the WNC alloys is found to be lower than that of WNF alloy. It is also observed that the micro-hardness of the W-grains are identical in all three alloys, while the micro-hardness of the matrix phase of WNF alloy (3.1 GPa) is slightly higher than that of WNC alloys (2.5–2.7 GPa). The difference in hardness, however, is much less pronounced than the difference in tensile strength and elongation values (Table 2) of WNF and WNCs. The tensile properties of WNF alloy are far superior to those of the WNC alloys. This might be due to the fact that W-grain size, and contiguity of WNF alloy are smaller as compared to those of WNC alloys [2] as shown in Table 2. The tensile fracture surfaces of both the WNC alloys and WNF specimen are shown in Fig. 3. While predominantly matrix or interface failure is observed in the case of WNCs alloys, the mode of failure is predominantly Wgrain cleavage fracture in WNF alloy at room temperature. Since the micro-hardness of the W grains in WNF and WNCs is similar, and the micro-hardness of the matrix phase in WNF is only slightly higher than that in WNC (Table 2), the vastly higher tensile
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Fig. 4. W, Ni, Fe and Cu elemental maps for (a) WNF, (b) WNC1 and (c) WNC2 alloy respectively.
strength of the WNF as compared to that of WNCs and the transgranular cleavage fracture in WNF possibly originated from finer W-grain size and better W/matrix phase cohesion and lesser contiguity (0.33) in WNF alloy as compared to those of WNC alloys. The other reported values [2,17–19] of microstructural parameters such as contiguity, W-grain size and mechanical properties are listed in Table 2 for comparison. The difference in contiguity between this study and literature is due to clustering of tungsten grains [1]. Where as the strength values are comparable, the difference in elongation values between this study and the literature is due to relatively higher contiguity and larger W-grain size of the WNF alloy produced in this study. WNC1 (95W–3.5Ni–1.5Cu) produced in this study showed inferior tensile properties (T.S. −660 MPa, El. −3%) as compared to 95W–5(Ni/Fe) alloy properties available in the literature [1,9,20]. It is interesting to note from the EPMA analysis that the matrix of the WNF alloy is richer in W (∼35 wt%) as compared to that of WNC alloys (∼20–23.4 wt%) (Table 4). W, Ni, Fe and Cu elemental maps for WNF, WNC1 and WNC2 alloy respectively are shown in Fig. 4. Some amount of porosity (marked by arrow in Fig. 3a and b)
was observed in the tensile fracture surfaces of WNC alloys but not in that of WNF alloy, which conformed well with the higher level of densification achieved in case of WNF alloy (Table 2). Lower tensile properties of WNC alloys might be partly due to the presence of these porosities. As WNF alloys show better room temperature tensile properties further room temperature tensile testing at different strain rates (0.1, 0.01, 0.001, 0.0001/s) were carried out and the results of these tests are displayed in Fig. 5. Analysis of these data show (Table 5) that in general at room temperature the yield stress and tensile stress of WNF increases and the % elongation diminishes with the increase in strain rate, which is expected. However, the mode of
Table 4 EPMA analysis (wt%) of the matrix phase of WHAs. Alloy
W
Ni
Cu
WNC1 WNC2 WNF
23.4 ± 1.9 20.2 ± 2.6 34.9 ± 6.5
58.8 ± 2.4 57.7 ± 1.4 49.9 ± 7.1
18.4 ± 2.5 23.7 ± 3.8
Fe
9.6 ± 1.4
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Table 5 Tensile properties of WNF at room temperature and at elevated temperatures vis-à-vis the room temperature and elevated temperature tensile properties of WHAs obtained from the literature. Alloy
Tensile test temperature(◦ C)
Strain rate (s−1 )
Yield stress (MPa)
WNF (tested in air)
25 25 25 25 500 700 25 500 700
0.1 0.01 0.001 0.0001 0.00003 0.000067 0.000083 0.000083 0.000083
922 803 716 630 368 370 ∼800 ∼450 ∼400
92.5W–5Ni–2.5Fe alloy [16] (tested partly under high vacuum)
Fig. 5. Tensile stress–strain curves of room temperature and elevated temperature tested samples which were prepared from sintered and heat treated WNF alloy.
fracture remained trans granular at all strain rates. The room temperature stress–strain curve at slow strain rate of 0.0001/s is wavy in nature; but this is absent at higher strain rates (Fig. 5). Hence, wavy nature in the tensile curve is not due to the presence of any porosity in the sintered and heat treated WNF alloy. In fact, no such microporosity was revealed by the micrograph in Fig. 3c. The wavy nature of the tensile curve indicates inhomogeneous deformation. The dislocation locking–unlocking phenomenon at low strain rate might be responsible for rendering the wavy tensile curve [21]. Room temperature tensile properties are comparable with available data in the literature [16]. The tensile properties of WNF alloy was also evaluated at elevated temperatures (773 K and 973 K), and the results are shown
Engineering stress at maximum load (MPa) 1054 1045 994 966 548 427 ∼1050 ∼700 ∼600
Total elongation (%) 25.4 26.8 37.5 35.1 4.90 3.00 ∼20 ∼15 ∼16
in Table 5 and Fig. 5. The results, when compared with room temperature test data (Table 5), show that both tensile strength and yield strength decreases by more than 40% at elevated temperatures, but a catastrophic decrease in the % elongation of the sample by more than 85% occurred upon rising the temperature. However, the catastrophic decrease in % elongation is not observed [16] at elevated temperatures (773 K and 973 K) even at a very slow strain rates (0.000083/s), when the WHA alloy specimens are tensile tested partly under high vacuum. Since 973 K is far lower than the recrystallization temperature of W (>1650 K [22]), the deterioration of the tensile properties at 973 or 773 K must be due to softening and embrittlement of the matrix phase by oxidation during the test conducted in air. Oxidation (Fig. 6b) on the fracture surface has occurred due to the fact that tensile test was carried out at a very slow strain rate in air. The fractograph of the high temperature tensile tested specimen is shown in Fig. 6a. Here the specimen has failed by predominantly matrix fracture during tensile testing, whereas at room temperature WNF failed predominantly by W-grain cleavage fracture (Fig. 3c). 4. Conclusions The tensile properties and hardness of WNF alloy at room temperature are superior due to the presence of finer W-grain size, lesser contiguity and porosity in its microstructure. WNF alloy fails under tension by W-grain cleavage fracture due to relatively stronger matrix phase and W/matrix bonding. The tensile properties and hardness of WNC alloys at room temperature are inferior due to the presence of coarse W-grain size, higher contiguity and porosity in their microstructure. WNC alloys fail under tension by matrix/interface fracture due to relatively weaker matrix phase and W/matrix bonding.
Fig. 6. (a) Fractograph of broken surface of WNF alloy which was tensile tested at 773 K and at a strain rate of 0.00003/s and (b) EDS graph obtained from the fracture surface.
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When WNF alloy is tested in air for prolonged time at elevated temperatures (773 and 973 K), its tensile strength and elongation remarkably deteriorates and the fracture changes to matrix failure mode due to weakening of the matrix phase. Acknowledgements Authors convey their sincere gratitude to Dr. G Malakondaiah, Distinguished Scientist and Director, DMRL, for giving permission to publish this work. The financial support provided by DRDO is thankfully acknowledged. They would like to thank the staff members of Powder Metallurgy Group, DMRL for their help in experimentation. Thanks are also due to Shri Bijoy Sarma, Division Head, Powder Metallurgy Division, for necessary support and useful discussion. One of the authors (JD) is indebted to Mr. A Saibabu, Sc. ‘F’, (retired) for his helpful suggestions and guidance to carry out this investigation. Authors gratefully acknowledge Shri A Sambasiva Rao for his support during EPMA work. References [1] R. Gero, L. Borukhin, I. Pikus, Mater. Sci. Eng. A 302 (2001) 162–167. [2] R.M. German, L.L. Bourguignon, Powder Metall. Def. Technol. 6 (1984) 117–131. [3] F.V. Lenel, Powder Metallurgy Principles and Application, MPIF, Princeton, 1980. [4] R.M. German, L.L. Bourguignon, B.H. Rabin, Powder Metall. 5 (1992) 3–13. [5] R. Gero, D. Chaiat, in: I. Minkoff (Ed.), High Density Tungsten Alloys, Mater. Eng. Conf., Israel, 1981, pp. 46–50.
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