UNCORRECTED PROOF. Matrix pools in a partially mechanically alloyed tungsten heavy alloy for localized shear deformation

Transcription

1 Materials Science and Engineering A000 (2001) Matrix pools in a partially mechanically alloyed tungsten heavy alloy for localized shear deformation Soon H. Hong a,hoj.ryu b, *, Woon H. Baek c a Department of Materials, Science and Engineering, Korea Ad anced Institute of Science and Technology, Guseong-dong, Yuseong-gu, Daejeon , South Korea b DUPIC, Korea Atomic Energy Research Institute, 150 Deokjin-dong, Yuseong-gu, Daejeon , South Korea c Agency for Defense De elopment, P.O. Box 35, Yuseong-gu, Daejeon , South Korea Abstract Received 16 July 2001; received in revised form 24 September 2001 A fabrication process for a tungsten heavy alloy with matrix pools was suggested in order to increase the susceptibility to localized shear deformation. A partial mechanical alloying process was introduced to form matrix pools intentionally within a tungsten heavy alloy. The partially mechanically alloyed powders were sintered by solid-state sintering, followed by liquid-phase sintering. The volume fraction of matrix pools was decreased with an increase in the secondary liquid-phase sintering time. Penetration test showed that the matrix pools were effective in enhancing the triggering of localized shear deformation Published by Elsevier Science B.V. Keywords: Tungsten heavy alloy; Mechanical alloying; Matrix pools; Localized shear deformation 1. Introduction Tungsten heavy alloys (WHAs) are used for high density structural applications, such as kinetic energy penetrators, counter weights, radiation shields and electrical contacts [1]. Liquid-phase sintering of mixed elemental powders of tungsten, nickel and iron at a temperature 1500 C has been a conventional fabrication process of WHAs [2]. There are two kinds of alloys used for kinetic energy penetrator applications. One is depleted uranium alloys, such as U Ti, U Mo and U Nb and the other is WHAs, such as W Ni Fe and W Ni Fe Co. Kinetic energy penetrators made of depleted uranium alloys exhibit superior penetration performance due to their so-called self sharpening behavior after impact against armor target, which results in a maintenance of small head diameter during penetration [3]. On the contrary, kinetic energy penetrators made of WHAs form mushroom-like head during penetration, which results in a poorer penetration performance compared to depleted uranium alloys [4]. It is needed to develop a high performance WHA to replace the depleted uranium alloy, since the uranium alloys cause the environmental contamination [5 8]. It has been known that the self sharpening behavior is related to localized shear deformation at high strain-rate deformation [9,10]. Recently, many researchers have introduced new processing techniques to modify microstructures of WHAs in order to enhance the localized shear deformation during penetration by the alloying with elements, such as Mo, Re and Mn [11,12], cyclic heat treatment [13], carburization of surface [14], solid-state sintering [15] and mechanical alloying [16,17]. Bose et al. showed that inhomogeneous distribution of matrix pools enhanced the initiation of adiabatic shear deformation, but they concluded that the reproducible fabrication of inhomogeneous distribution of matrix pools was impossible [18]. In this study, the partial mechanical alloying process, which consists of mechanical alloying of powders with a composition of matrix phase and mixing with the remaining W powders, is investigated to form intentional formation of matrix pools in a WHA. Formation * Corresponding author. address: hjryu@kaeri.re.kr (H.J. Ryu) /01/$ - see front matter 2001 Published by Elsevier Science B.V. PII: S (01)

2 2 of localized shear deformation in a WHA rod during penetration against steel targets is observed. S.H. Hong et al. / Materials Science and Engineering A000 (2001) Experimental procedures Blended powders of tungsten, nickel and iron with a composition of 30W 56Ni 14Fe were mechanically alloyed in a Spex mill for 1 h and followed by blending with W powders by turbular mixer to form final composition of 93W 5.6Ni 1.4Fe. The mechanically alloyed and blended powders were consolidated into green compacts by die compaction under 100 MPa. The green compacts were solid-state sintered at 1300 C in a hydrogen atmosphere. The solid-state sintered WHA was subsequently liquid-phase sintered, again at a temperature of 1470 C, with varying sintering time from 4 to 90 min. The two-step sintered specimens were annealed at 1150 C for 1 h in a nitrogen atmosphere, followed by water quenching to prevent hydrogen embrittlement and impurity segregation during the cooling stage. Sizes of tungsten particles, volume fractions of matrices and W/W contiguities in two-step sintered WHAs were characterized from microstructural analysis using scanning electron microscopy. Yield strength Fig. 2. The scanning electron micrographs showing the microstructures of 93W 5.6Ni 1.4Fe tungsten heavy alloys after solid-state sintering at 1300 C of (a) blended powders of mechanical alloyed 30W 56Ni 14Fe powders and additional W powders to form 93W 5.6Ni 1.4Fe and (b) fully mechanically alloyed 93W 5.6Ni 1.4Fe powders. Fig. 1. The scanning electron micrographs of 30W 56Ni 14Fe powders of (a) mechanically alloyed for 5 min and (b) mechanically alloyed for 60 min. and elongation were characterized using tensile tests of specimens with gauge length of 25 mm at a strain-rate of s 1. A penetration test of WHA rods against mild steel with a thickness of 80 mm was performed at a striking velocity of 1100 m s Results and discussion Fig. 1 shows the microstructure of mechanically alloyed 30W 56Ni 14Fe powders using Spex mill. 30W 56Ni 14Fe is similar to the composition of solid solution matrix of 93W 5.6Ni 1.4Fe WHA [13]. 30W 56Ni 14Fe powders mechanically alloyed for 5 min show a mixture of bright phase of tungsten and dark phase of nickel and iron in a coarse powder formed due to the repeated welding of elemental powders between colliding balls in Spex mill, as shown in Fig. 1(a). Fig. 1(b) exhibits a homogeneous structure of

3 S.H. Hong et al. / Materials Science and Engineering A000 (2001) Fig. 3. The variation of microstructure of partially mechanically alloyed tungsten heavy alloy after liquid-phase sintering for (a) 4 min, (b) 8 min, (c) 15 min and (d) 90 min at 1470 C. mechanically alloyed 30W 56Ni 14Fe powders after mechanical alloying of 60 min in Spex mill. 30W 56Ni 14Fe powders mechanically alloyed for 60 min were mixed with tungsten powders with a ratio of 1:9 to fabricate 93W 5.6Ni 1.4Fe WHA and this process was named as a partial mechanical alloying process to discriminate from a full mechanical alloying process, where whole 93W 5.6Ni 1.4Fe powders are mechanically alloyed at the same time [15]. Fig. 2(a) shows a microstructure of a partially mechanically alloyed WHA after solid-state sintering at 1300 C. Inhomogeneous distribution of matrix pools after solid-state sintering of the partially mechanically alloyed WHA is outstandingly different from that of fully mechanically alloyed WHA, as shown in Fig. 2(b) [15]. The partial mechanical alloying process was found effective in forming matrix pools intentionally in a WHA. The size of matrix pools is similar to the powder size after mechanical alloying of matrix composition powders. The matrix pools are thought to result from the rapid sintering of mechanically alloyed powders. The solid-state sintered WHAs were subsequently sintered at 1470 C, in order to control the volume fraction of matrix pools and W/W contiguity, which is the area fraction of contiguous W/W interface. This process is as two-step sintering, in order to form a fine microstructure of mechanically alloyed WHAs [19]. When sintered at 1470 C, the tungsten particles were spherodized due to the formation of the liquid-phase matrix. The matrix volume fraction increased and W/W contiguity decreased due to the increase in tungsten Fig. 4. The variation of vol.% of matrix pools and W/W contiguity of a partially mechanically alloyed tungsten heavy alloy with increasing secondary liquid-phase sintering time.

4 4 S.H. Hong et al. / Materials Science and Engineering A000 (2001) Fig. 3 shows the evolution of microstructures of WHAs with matrix pools after liquid-phase sintering at 1470 C with a varying sintering time from 4 to 90 min. The average size of tungsten particles increased with increasing liquid-phase sintering time, according to the diffusion controlled coarsening mechanism [21]. It was observed that the matrix pools disappeared when the sintering time increased due to the redistribution of the matrix phase between tungsten particles by the effect of capillary force [22]. W/W contiguity decreased with increasing liquid-phase sintering time due to a decrease in the volume fraction of the matrix pool, as shown in Fig. 4. Mechanical properties of two-step sintered WHA Fig. 5. The variation of yield strength of elongation of partially mechanically alloyed tungsten heavy alloy with increasing secondary liquid-phase sintering time. Fig. 6. The fractography of a partially mechanically alloyed tungsten heavy alloy after liquid-phase sintering for (a) 8 min and (b) 90 min at 1470 C. Fig. 7. The characteristic microstructure of localized shear deformation in a tungsten heavy alloy after a penetration test. solubility in matrix phase and the decrease in a dihedral angle of tungsten/matrix interface after secondary liquid-phase sintering [20]. Fig. 8. The micrograph showing the localized shear deformation in residual tungsten heavy alloy penetrator after penetration test against mild steel at a striking velocity of 1100 m s 1.

5 S.H. Hong et al. / Materials Science and Engineering A000 (2001) were evaluated by the tensile tests, as shown in Fig. 5. Yield strength decreased and elongation to failure increased with increasing secondary liquid-phase sintering time. Ryu et al. showed that yield strength of a WHA is correlated with microstructural parameters as follows [17,23] y = 0 +k 1 V M 1/2 (1) DV M where y is the yield strength, 0 is intrinsic strength, k is constant, V M is the matrix volume fraction and D is the average tungsten particle size. According to Eq. (1), yield strength of WHA increased with decreasing tungsten particle size and matrix volume fraction. An increase in elongation to failure with increasing secondary sintering time was associated with decreased W/W contiguity. Fig. 6 shows micrographs of the fractured surface of WHAs liquid-phase sintered for 5 and 60 min. The fractured surface of a WHA liquidphase sintered for 5 min consisted of tungsten/tungsten boundaries, while tungsten cleavage was predominantly shown in a WHA liquid phase sintered for 60 min. Because the tungsten/tungsten boundary are the weakest boundary [24], a larger tungsten/tungsten contiguity resulted in a reduced elongation to failure. The characteristic localized shear deformation in a WHA is shown in Fig. 7 [25]. The self-sharpening behavior of kinetic energy penetrators is known to occur through the propagation of cracks along the heavily deformed localized shear deformation zone. Fig. 8 shows the resulting microstructure of WHA after penetration test against mild steel target at a striking velocity of 1100 m s 1. Localized shear bands, triggered at matrix pool, were observed from the residual penetrator of WHA. It is generally understood that shear deformation is concentrated within the ductile matrix phase [26]. Zhou and Clifton [27] showed that the matrix phase in WHA served as a perturbation to deformation that enhanced the initiation and development of shear bands and Stevens and Batra [28] reported that softer matrix introduced sufficient perturbation to incite localization of deformation. Therefore, formation of matrix pools was found to be effective in enhancing the initiation of localized shear deformation in WHA. The partial mechanical alloying, suggested in this study, can be a method of forming the initiation sites of localized shear deformation. Because the improvement of the actual performance of the penetrator is the most important purpose of the microstructural control of tungsten heavy alloys including this study, more penetration tests with varying size and volume fraction of the matrix pool in a tungsten heavy alloy are needed to clarify the relationship between the penetration depth and the size and volume fraction of the matrix pool. 4. Conclusions A kind of mechanical alloying process is suggested to form inhomogeneous distribution of solid solution matrix phase in a WHA. A 93W 5.6Ni 1.4Fe WHA with matrix pools was fabricated by two-step sintering of WHA powders, in which partially mechanically alloyed matrix composition powders were added. Mechanical properties of sintered WHA were dependent on microstructural parameters, such as the average size of a tungsten particle and W/W contiguity. Intentional formation of matrix pools was found effective in enhancing the triggering of localized shear deformation in WHA. References [1] W.D. Cai, Y. Li, R.J. Dowding, F.A. Mohamed, E.J. Lavernia, Rev. Particul. Mater. 3 (1995) 71. [2] R.M. German, in: A. Bose, R.J. Dowding (Eds.), Proceedings of the International Conference on Tungsten and Tungsten Alloys, MPIF, New Jersey, 1992, p. 3. [3] S.P. Andrew, R.D. Caligiuri, L.E. Eiselstein, in: A. Crowson, E.S. Chen (Eds.), Tungsten and Tungsten Alloys Recent Advances, TMS, Warendale, 1991, p [4] L.S. Magness, T.G. Farrand, Proceedings of Army Science Conference, Durham, 1990, p [5] R.J. Dowding, M.C. Hogwood, L. Wong, R.L. Woodward, in: A. Bose, R.J. Dowding (Eds.), Proceedings of the International Conference on Tungsten and Refractory Metals, MPIF, New Jersey, 1994, p. 3. [6] D.K. Kim, S. Lee, H.S. Song, Met. Mater. Trans. A29 (1998) [7] D.K. Kim, S. Lee, H.J. Ryu, S.H. Hong, J.W. Noh, Met. Mater. Trans. A31 (2000) [8] S. Pappu, C. Kennedy, L.E. Murr, L.S. Magness, D. Kapoor, Mater. Sci. Eng. A262 (1999) 115. [9] K.T. Ramesh, R.S. Coates, Met. Trans. A23 (1992) [10] A. Bose, H.R.A. Couque, J. Langford, Int. J. Powder Met. 28 (1992) 383. [11] A. Bose, G. Jerman, R.M. German, Powder Met. Inter. 21 (1989) 9. [12] E.P. Kim, M.H. Hong, W.H. Baek, I.H. Moon, Met. Mater. Trans. A30 (1999) 627. [13] J.W. Noh, E.P. Kim, H.S. Song, W.H. Baek, K.S. Churn, S.J.L. Kang, Met. Trans. A24 (1993) [14] S.W. Jung, D.K. Kim, S. Lee, J.W. Noh, S.J.L. Kang, Met. Mater. Trans. A30 (1999) [15] H.J. Ryu, S.H. Hong, W.H. Baek, Mater. Sci. Eng. A291 (2000) 91. [16] M.L. Ovecoglu, B. Ozkal, C. Suryanarayana, J. Mater. Res. 11 (1996) [17] H.J. Ryu, S.H. Hong, W.H. Baek, J. Mater. Process. Tech. 63 (1997) 292. [18] A. Bose, H. Coque, J. Lankford Jr, in: A. Bose, R.J. Dowding (Eds.), Proceedings of the International Conference on Tungsten and Tungsten Alloys, MPIF, New Jersey, 1992, p [19] H.J. Ryu, S.H. Hong, S. Lee, W.H. Baek, Met. Mater. 5 (1999) 185. [20] B.H. Rabin, A. Bose, R.M. German, Int. J. Powder Metall. 25 (1989) 21.

6 6 S.H. Hong et al. / Materials Science and Engineering A000 (2001) [21] R.M. German, E.A. Olevsky, Met. Mater. Trans. A29 (1998) [22] H. Fredriksson, A. Eliasson, L. Ekbom, Int. J. Refract. Met. Hard Mater. 13 (1995) 173. [23] M.F. Ashby, Phil. Mag. 18 (1970) 399. [24] K.S. Churn, D.N. Yoon, Powder Met. 22 (1979) 175. [25] H.J. Ryu, Ph.D. Thesis, Korea Advanced Institute of Science and Technology, [26] L. Ekbom, Mod. Dev. Powder Met. 14 (1980) 177. [27] M. Zhou, R.J. Clifton, J. Appl. Mech. 64 (1997) 487. [28] J.B. Stevens, R.C. Batra, Inter. J. Plast. 14 (1998) 841.