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Microstructure and mechanical properties of W-Zr reactive materials
Materials Science & Engineering A 660 (2016) 205–212
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure and mechanical properties of W-Zr reactive materials Huilan Ren, Xiaojun Liu, Jianguo Ning n State Key Laboratory of Explosion Science and Technology, Beijing Institute of Technology, Beijing 100081, PR China
art ic l e i nf o
a b s t r a c t
Article history: Received 20 November 2015 Received in revised form 13 January 2016 Accepted 4 February 2016
Three different batches of tungsten-zirconium (W-Zr) reactive material were prepared by hot-pressing elemental powder mixtures. The first sample had a Zr:W mass ratio of 66:34. The second used zirconium hydride (ZrH2) with a (ZrH2):W mass ratio of 66:34. The third had a Zr:W mass ratio 43:57. These batches were numbered #1 to #3. To investigate the mechanical properties of the W-Zr reactive materials, samples were subjected to different strain rates using a materials testing machine and a modified split Hopkinson pressure bar (SHPB). The quasi-static compressive strengths of the three sample batches all exceeded 1022 MPa, with batch #1 sample having the highest value at about 1880 MPa. This could be ascribed to the sample exhibiting a more transgranular and dimpled fracture in the W2Zr intermetallic phase, as demonstrated by the microstructure of the fracture surface, observed by scanning electron microscopy (SEM) using an energy-dispersive spectrometer (EDS). To perform a constant strain rate experiment on this high-strength but brittle material, ramp loading using copper sheet was adopted in this study. All of the samples produced strong, bright flames when subjected to shock loading and exhibited a high compressive strength of approximately 1060 to 2690 MPa. High-resolution X-ray diffraction (XRD) was performed on the original samples and residues after the SHPB test, showing that the Zr and ZrC phases of the batch #1 and batch #3 samples, and the ZrC0.32H1.2 phase of the batch #2 sample are the active components in the reaction with air. Some small balls of ZrO2 reaction product were found not to exhibit any crystalline tungsten oxide on the residue surface. These results suggest that the batch #1 reactive material has huge potential to take the place of inert steel because of its high strength and high energy level, as well as having a density close to that of steel. & 2016 Elsevier B.V. All rights reserved.
Keywords: W-Zr Reactive materials Mechanical properties Microstructure Fracture
1. Introduction Reactive materials are a special class of energetic materials with a high energy density, which react with themselves or combust violently with air and rapidly generate high pressures and temperatures when a certain reaction initiation energy level is achieved, leading to additional damage. The reactive materials can be adapted to a variety of applications such as target damage through the addition of structural reactives, propellant/explosive additives, and manufacturing. Typical reactive materials usually consist of two or more solid-state reactants, unlike explosives that are highly sensitive to ignition, that together form a thermo-chemical mixture. These materials mainly consist of thermites, intermetallics, metal/fluorine systems, metastable intermolecular composites (MIC), nano-laminates, and metal hydrides [1–3]. Among these, the intermetallic-forming reactive materials are not only energy reactants, but can also act as structural components due to their good mechanical properties. Due to their unique n
Corresponding author. E-mail address: [email protected] (J. Ning).
http://dx.doi.org/10.1016/j.msea.2016.02.009 0921-5093/& 2016 Elsevier B.V. All rights reserved.
characteristics, the applications of intermetallics have recently diversified. They have been applied to the likes of missiles, given that their strength allows them to be used as a portion of the missile casing. The energy stored within the material will be released when subjected to the loading of impact and thus result in an increase in the explosive power of the missile, relative to that of a design using a traditional steel casing [4–6]. Intermetallics are an important subclass of reactive materials that have been studied by numerous researchers over the last several decades. They include Ni-Al and Al-W powder systems, and their study has included their initiation and characterization. The ultrafast reactions of Ni-Al powders when subjected to shock loading were first observed by Bennett et al. [7]. They found that the particle size and morphology had a significant influence on the shock ignition behavior of the Ni-Al powder mixtures in plate impact tests. This was subsequently proven by Eakins and Thadhani [8,9]. A modified Asay shear impact experiment showed that the mechanical dry milling time of the Ni-Al composites could significantly affect the microstructure and the mechanical impact ignition threshold [10]. The results of numerical simulations showed that the configuration of the Ni-Al powders played an important role in the coupled thermal–mechanical behavior under
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shock wave loading [11]. The dynamic behavior and fracturing of Al-W granular composites have been investigated and the results revealed that the morphology of the W particles had a significant effect on the dynamic behavior [12–14]. Explosively driven fragmentation experiments were performed on Al-W rings. Furthermore, the mesoscale fragmentation mechanisms that could give rise to enhanced pulverization were studied by Olney et al. [15– 17]. The failure mechanisms of axial splitting and shear failure were observed in the continuous phase of different explosively consolidated metal-aluminum powder mixtures in response to an impact [18,19]. The research results indicated that the W-Zr fragment had a capacity in excess of 93 W when initiating a charge and fuel tank, resulting from the large scattering angle and multifire point ignition. The initiation by W-Zr fragment has mainly a mechanics, while the other case has mainly a shock wave [20,21]. Impact-initiated experiments with a quasi-sealed test chamber were conducted with a W-Zr alloy and the results showed that the impact velocity could seriously impact the reaction behavior [22,23]. The W particles in W-Zr mechanical alloys that were normally regarded as being inert were proven to be involved in the reaction and released large amounts of energy under explosive tests [24,25]. Clouds of Zr dust with an average particle size of about 3 mm were susceptible to spontaneous ignition in air at room temperature. The affinity of Zr with oxygen was great, and it could also ignite in a carbon dioxide atmosphere at approximately 621 °C and in nitrogen at approximately 788 °C [26,27]. The dynamic compression mechanical properties of a W-Zr alloy under different dynamic loads were investigated and the results indicated that Zr flames reach very high temperatures, while W-Zr composites still retain their desirable structural properties before ultimately reacting [22,28]. In this study, to further improve the energy level and mechanical properties of reactive materials, hot-pressing techniques were adopted to prepare three different batches of W-Zr reactive material, the density of which was the same as steel, which could be applied to reactive material fragmentation weapons with the same kinetic energy. We investigated the differences between the microstructures and mechanical properties of the three batches of W-Zr reactive materials, when subjected to a range of loading rates.
2. Experimental setup The W-Zr reactive material was prepared from an elemental W powder (purity 99.9%, average particle size 3 mm), Zr (purity 99%, average particle size 37 mm) and ZrH2 (purity 99%, average particle size 3 mm) using hot-pressing. The powders were mixed in three different ways. The batch #1 and #2 samples had identical mass ratios with 34% of the W component. The batch #1 and #3 samples used the same Zr powder but at different mass ratios. The compositions and proportions are listed in Table 1.The preparation of the W-Zr reactive materials involved mixing and sintering. The process is shown in Fig. 1. During the sample preparation, the powders were mixed with absolute ethyl alcohol in a low-speed V-shaped blender for 6 h and then dried at 60 °C in an electrically heated forced-air drum-type
Tungsten Powder
Mixing
Zirconium/Zirconium Hydride Powder
Dry
Mixed Powder Ar Furnace Cooling W-Zr
Sintered
Fig. 1. W-Zr preparation process using hot-pressing.
Fig. 2. Typical hot-pressing curves for W-Zr.
oven for 24 h. Then, the dried mixtures were screened through a 200-mesh sieve and placed in a graphite die (Φ65 mm). Fig. 2 shows the temperature and pressure history in the sintering cycle, which was as follows: The pressure was increased to 20 MPa over 2 h and then held constant for 3.5 h. Thereafter, the pressure was reduced to 0 MPa at a rate of 60 MPa/h. The oven temperature was then increased to 1000 °C over 1 h, then slowly raised to 1500 °C at a rate of 500 °C/h and then held at 1500 °C for 3 h. After that, the temperature was decreased to 200 °C at a rate of 600 °C/h and then allowed to naturally cool to room temperature in the furnace. The entire sintering process was conducted in an argon atmosphere. For this study, each batch of the W-Zr reactive material was formed into two different sizes of sample, specifically, Φ10 mm 10 mm and Φ5 mm 5 mm, as shown in Fig. 3. Quasi-static compression tests were performed using a materials testing machine (Kexin WDW-300, China) for the samples of the three batches at room temperature, and dynamic tests were conducted using a split Hopkinson pressure bar for the W-Zr reactive material samples measuring Φ5 mm 5 mm. The apparatus consisted of a 200-mm striker bar, a 600-mm incident bar, and a 600-mm transmitting bar. All three bars were made of maraging steel with a diameter of 14.5 mm. Strain gauges were bonded to the midpoints of the incident bar and transmitted bar to record the stress waves propagating along the bars. Two high-strength
Table 1 Composition and proportion. Batch
Zr (37 mm)
ZrH2 (3 mm)
W (3 mm)
#1 #2 #3
66 0 43
0 66 0
34 34 57
1500oC Pressure
Fig. 3. Initial samples of W-Zr reactive material.
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Fig. 4. Modified SHPB apparatus.
steel spacers with the same generalized wave impedance as the bar were added between the specimen and the incident and transmission bars to protect the end surfaces against damage. The layout of the apparatus is shown in Fig. 4.
3. Results and discussion 3.1. Initial microstructures The microstructures of the samples were characterized by SEM with EDS (Hitachi S-4800, Japan). The obtained images are shown in Fig. 5. X-ray diffraction (Bruker D8 Advance, Germany) analyses were performed to identify the phases in the reactive materials. The results are shown in Fig. 6. The densities and theoretical maximum specific density (%TMD) of the three batches are listed in Table 2. As can be seen in Fig. 5, there are significant differences between the microstructures of the three batches, which illustrates the fact that the components are also different. Fig. 6 shows that the samples mainly consist of ZrC, W2Zr, and Zr phases in batch #1, ZrC0.32H1.2, W, and Zr phases in batch #2, and W, W2Zr, and ZrC phases in batch #3. The micrograph of the batch #1 sample shows the presence of only two phases (Fig. 5(a)). The dark regions are the Zr and ZrC phases, while the white regions are the W-Zr intermetallic component, which was detected by EDS (Fig. 7). Due to the high temperature, the graphite paper used to ensure the easy release of the mold was involved in the reaction to form a ZrC phase, which can be seen in the Zr-C phase diagram (Fig. 8(a)). The phase diagram for the W-Zr system indicates that, in that region with more than 20 wt% Zr and in which the temperature reaches between 1400 and 2210 °C, there is one intermetallic phase, that is, W2Zr (Fig. 8(b)). Therefore, the composition of the bright phase in batch #1 corresponds to W2Zr. The intermetallic phase is distributed evenly in the ZrC and Zr phases, although with a variety of shapes and sizes. The density of the batch #1 sample is over 99% TMD, which is close to the fully dense material. In the batch #2 sample, many of the irregular W particles were found to be surrounded by a continuous gray matrix, as shown in Fig. 5(b). The
Fig. 6. XRD patterns of samples: (a) #1, (b) #2, (c) #3. Table 2 Density and %TMD of W-Zr reactive Materials of three batches. Batch
Density (g/cm3)
% TMD
#1 #2 #3
8.34 8.01 9.15
99.2% 95.3% 87.5%
ZrH2 can remain stable after long-term aging at room temperature, since it quickly reacts with the oxygen in the air and forms a nanometer-thin layer of oxide that prevents further oxygen diffusion into the material. Nevertheless, this protective effect is lost when the temperature exceeds 400 °C and the change in the composition due to the oxidation can usually be neglected. Norton proposed that ZrC can be produced by heating a mixture of ZrH2 and graphite to 1800 to 2200 °C, and pointed out that long-term heating has a significant effect on the decomposition of the ZrH2
Fig. 5. SEM micrographs of samples before test: (a) #1, (b) #2, (c) #3.
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Table 3 Quasi-static compressive results for three batches of W-Zr reactive Materials.
Fig. 7. EDS result for white region of batch #1 sample.
hydrides [29]. In this experiment, however, the ZrH2 and W powders were sintered in a graphite die in an Ar gas atmosphere rather than in a vacuum. Because of the lower temperature (1500 °C) and the decomposition of a small amount of carbon, no ZrC was formed with the batch #2 sample. The gray phase ultimately proved to be ZrC0.32H1.2 and Zr that transformed from ZrH2 particles in the sintering process. Given the high level of activity of the small particles, the remaining Zr was easily agglomerated. The density of the batch #2 sample was 95.3% TMD, slightly less than that of batch #1. Relative to batch #1, batch #3 exhibited a smaller white region with many evenly distributed holes (Fig. 5(c)). The density of the sample is thus highly dependent on the sintering temperature. A sintering temperature of 1500 °C is too low for the batch #3 sample with a higher volume fraction of W. As a result, the grain growth occurs slowly and the crystal growth is incomplete, which leads to there being a large number of pores and thus the density is lower (87.5% TMD). 3.2. Fracture observation The quasi-static test results indicate that all three batches of the W-Zr reactive materials are a typical brittle material with a high level of strength that exhibits a good linear elastic stressstrain curve under quasi-static loading. The compressive strength
Batch
Strain (%)
Stress (MPa)
#1 #2 #3
1.0 0.7 0.4
1880 1200 1022
of the batch #1 sample is higher than that of the other two batches. The details are listed in Table 3. Scanning electron microscopy was used to investigate the failure mechanisms and the typical fracture surface morphologies of the samples after different levels of plastic deformation. The resulting images are shown in Fig. 9. A zigzag crack appeared in the fracture surface of the batch #1 sample, as shown in Fig. 9(a). It can be seen that the fracture surface consists of two different zones. The white circle indicates where the Zr phase exhibits a transgranular fracture. From the top view of the fracture (Fig. 9(b)), the detailed fracture morphology exhibits a few dimples with features of plastic fracture in region A and a predominantly cleavage river pattern in region B. Fig. 9(c) shows that the W2Zr intermetallic presents an aggregated and irregular polyhedron appearance with particle sizes ranging from 1 mm to 3 mm. An enlarged view of the crack propagation path in the W2Zr phase is shown in Fig. 9(d). As indicated by the arrows, cracks appeared in the crystal grains and along the grain boundaries, implying that these interactions absorb the energy of the crack propagation during the fracture process, and that the fracture mode is a mix of transgranular and intergranular. These results imply that transgranular fractures occurred in the Zr and ZrC phase, while a mixture of transgranular, intergranular, and dimple fractures occurred in the W2Zr phase, thus leading to a toughening of the reactive material and the attainment of the highest compressive strength. The microstructure of the batch #2 sample surface is shown in Fig. 9(e). Here, ZrC0.32H1.2 phase transgranular fractures can be easily observed. However, the microstructure of the shear fracture was very different from that of the tensile fracture, and a large number of ZrC0.32H1.2 fragments, with the appearance of wood chips with many burrs, can be observed on the surface, as shown in Fig. 9(f). Fig. 9(g) shows transgranular fractures in the ZrC0.32H1.2 phase and many spherical-shaped W particles with sizes of 1–3 mm in the fracture surface, as indicated by the circles. It is evident that the integrated W particles were separated from the matrix as a result of the intergranular fractures during the test. After the sintering and the reaction that generated the ZrC0.32H1.2, there was some Zr residue and transgranular fractures can be observed in the tensile fracture region (see Fig. 9(h)). The experiments showed
Fig. 8. Phase diagram: (a) Zr-C system, (b)W-Zr system.
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Fig. 9. SEM images of fracture surfaces after quasi-static tests: (a)-(d) batch #1, (e)-(h) batch #2, (i)-(l) batch #3.
that transgranular fractures with a small intergranular fraction were a characteristic of the batch #2 reactive material. Fig. 9(i) shows that the batch #3 sample had a very loose structure, and the cracks progress along the grain boundaries by connecting with other microcracks or across the grains as intergranular fractures. Some Zr particles exhibited a serious deformation in the shear fracture region, which is clearly demonstrated in Fig. 9(j). For clarification, a magnification of Fig. 9(i) is shown in Fig. 9(k). The fracture surface exhibited intergranular cracks and some W2Zr can be observed. Relative to batch #2 (Fig. 9(f)), the fracture surface of the batch #3 samples exhibits fewer cleavage planes, but more holes, as shown in Fig. 9(l).
3.3. Dynamic behavior In a dynamic test to determine the material properties, the sample is required to deform nearly uniformly at a constant strain rate under dynamically equal stress. For brittle materials that are loaded dynamically with the SHPB apparatus, ramp loading is typically employed to facilitate a constant strain rate experiment [30]. The typical original waveforms of the SHPB test with a pulse shaper (copper sheet) are shown in Fig. 10(a). The pulse shaping has an obvious effect on the elimination of the high-frequency signals and the delay in the incident pulse [30]. The incident wave
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Fig. 10. Results of pulse-shaped SHPB experiment with batch #1: (a) Strain-time signals, (b) Interface stresses.
becomes less steep and gives rise to ramp loading. For the batch #1 sample, tested with a gas chamber pressure of 0.1 MPa, for example, the rise time (from point A) of the incident wave was about 37 μs. For a time between 201 μs (point C) and 210 μs, the reflected wave signal was approximately equal to a constant value. Thereafter (point D), the reflected signal presented a steep rise, indicating that the sample has already been damaged or has failed. The strain signal of the sample gradually changed from a gentle slope to a steep slope between 138 μs (point B) and 150 μs. Its value changes with time, almost as a straight line, up until 162 μs, which shows that the sample deformed at a nearly constant strain rate during this period. The elastic wave velocity of the sample is about 4800 m/s and the characteristic time t¼ 1 μs (t is the duration of the stress wave required to propagate from one side in the sample to another side). Based on the literature, the rise time of the incident wave should be greater than 8t, which is a prerequisite condition for achieving the dynamic stress equilibrium in the brittle sample [31]. In this work, the stress σ1 at the incident bar/sample interface and the stress σ 2 at transmission bar/sample interface as obtained from a pulse-shaped SHPB experiment were as shown in Fig. 10(b). It can be seen that, after 10 μs, these interface stresses are in reasonably good agreement. This suggests that the samples were nearly in a state of dynamic stress equilibrium over most of the duration of the test. The sample deforms at a nearly constant strain rate under dynamically equilibrated stresses, which indicates that the data obtained from the pulse-shaped SHPB experiment was valid, and that the stresses and strains in the
samples could be calculated as follows:
⎧ 2C εR (t ) ⎪ ε ̇ (t ) = L0 ⎪ t ⎪ 2C ⎨ ε (t ) = [εR (t )] dt L0 0 ⎪ ⎪ A ⎪ σ (t ) = − Eε T (t ) A0 ⎩
∫
(1)
In Eq. (1), E, A, and C are the Young's modulus, cross-sectional area, and elastic bar wave velocity of the bar, respectively, and A0 and L0 are the cross-sectional area and the original length of the specimen, respectively. The reactive materials' maximum stress-strain relationships under quasi-static and high strain rates ranged from 200 to 1200 s 1. These are shown in Fig. 11(a). They accurately describe the compressive mechanical behaviors of the W-Zr reactive material. The quasi-static and dynamic stresses are plotted as a logarithmic function of the strain rate in Fig. 11(b). The results show that the compressive stress of all three batches of reactive materials, as determined by the dynamic tests, was higher than in the quasi-static tests, due to in-situ densification under dynamic conditions exceeding the fracture kinetics. The batch #1 samples had the highest compressive strength, and a similar behavior was also observed in the quasi-static tests. The best performance was obtained with the batch #1 samples, which can be explained by the fact that the intermetallic phases of the W2Zr provided a higher interface strength. Strain hardening
Fig. 11. Compression results for three batches under different strain rates: (a) Maximum stress vs. failure strain, (b) Maximum stress vs. strain rate.
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Fig. 12. Violent reaction of batch #1 sample when subjected to strong impact.
behavior under dynamic loading was observed, with the value ranging from 0.01 to 0.015. As shown in Fig. 11(a) (black squares), the stress-stain relationship followed a straight line. In the same way as with the trends with the batch #1 reactive material, the compressive strength of the batch #2 samples was higher at the high strain rate. The batch #2 samples had the same W mass ratio as batch #1, but were more brittle. However, the batch #3 samples had the weakest compressive strength and strain, which was attributed to the structure having many holes. The compressive strength at a strain rate of 470 s 1 differed slightly from that of the static condition (1060 MPa and 1022 MPa, respectively). The failure strain was 0.012 at a strain rate of 1200 s 1, which was more than three times that at a strain rate of 470 s 1. The samples failed, reacted, and rapidly released large amounts of chemical energy in the form of heat, light, and pressure when subjected to large impact loads. The reactions are shown in Fig. 12. This shows that the material has a very high energy content, with batch #1 exhibiting the greatest content, followed by batch #2, and then batch #3. This was determined from the size of the generated flame. According to the results of the XRD spectrum observations, the phase compositions of many of the residues are the same as those of the original samples, as is evident in Figs. 6 and 13. In addition, new diffraction peaks of the ZrO2 phases can be seen in the XRD pattern, which suggests that the Zr, ZrC, and ZrC0.32H1.2 phases are the active components. After the application of a strong impact, the high temperature and surface tension gave rise to a balling effect of the reaction product. As can be seen in Fig. 14, many small spherical balls with particle sizes of 1–10 mm are present in the
Fig. 13. XRD pattern of investigated W-Zr composite after SHPB test:(a) batch #1, (b) batch #2, (c) batch #3.
Fig. 14. SEM images of fracture surfaces of batch #1 sample after SHPB test.
fracture surface morphologies of the batch #1 sample after the SHPB test, and were finally confirmed to be Zr oxide through an EDS analysis. It should be noted that the Zr has a great affinity for oxygen, especially when the powder is very fine [27]. The ignition characteristics of the Zr and ZrH2 particles in different atmospheres were investigated by Cooper [26]. These results revealed that layers of Zr dust with particle diameters of less than 45 mm could be ignited at 530 °C in a nitrogen atmosphere, but ZrH2 powder with a particle diameter of less than 4.7 mm could not be ignited in a nitrogen atmosphere, even at 850 °C. This showed that, even though the ZrH2 powder had a smaller particle size, its affinity for nitrogen was weaker than that of the Zr powder. In addition, kinetic studies of the oxidation of ZrC were performed at about 600 °C, resulting in the formation of the ZrO2 phase [32]. The batch #1 and batch #2 samples had a higher mass ratio of elemental Zr compared to the batch #3 sample, while the batch #1 sample contained much more “fuel,” such as the Zr component, than the another two batches. These results also demonstrate why the batch #1 sample exhibited the highest luminosity during the SHPB test. Upon applying an impact load to a metal or alloy with a high stacking fault energy, due to the high strain rate and large deformation, the distribution of the dislocations becomes more uniform and the dislocation density increases, resulting in the grain refining effect being better and making it easier to produce nano-structural characteristics in the material [33]. Severe plastic deformation of the Zr particles, as well as a few clusters of irregular particles with average diameters of about 100 nm, can be observed in the batch #1 sample after the SHPB test, which can clearly be confirmed from the high-magnification SEM image (Fig. 15). These nano-structures appeared primarily near the
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Fig. 15. SEM images of configuration of batch #1 residue at higher magnification and nano-structure phase formed.
interface between the Zr phase and W2Zr phase, and this nano material was ultimately proven to be W2Zr as a result of EDS identification.
4. Conclusions
1. Three different batches of W-Zr reactive materials were successfully prepared using a hot-pressing process. The density of the batch #1 and batch #2 samples was in excess of 95% TMD, while that of the batch #3 sample was only about 87.5% TMD since the sintering temperature of 1500 °C was too low with a higher volume fraction of W. 2. The quasi-static compressive strengths of the three batches were higher than 1022 MPa, with the batch #1 sample having the highest value of about 1880 MPa. The microstructures of the fracture surfaces were investigated by SEM and EDS. The results showed that the predominant fracture modes of the W-Zr reactive materials were transgranular and intergranular, while exhibiting brittleness damage, while transgranular and dimple fractures were more predominant in the W2Zr intermetallic phase, leading to batch #1 sample toughening. 3. Waveform shaping technology was used to deform the samples at a nearly constant strain rate under dynamically equilibrated stresses, since the W-Zr reactive materials feature high strength but are brittle. A strain gauge was directly attached to the sample surface to acquire more accurate data, and the dynamic compressive strength was about 1060–2690 MPa. 4. All of the samples produced a clear flame when subjected to shock loading and an XRD analysis showed that the Zr, ZrC, and ZrC0.32H1.2 phases were the active components in the reaction with air, while reaction products of ZrO2 could be found in the residue surfaces. The batch #1 reactive material offers excellent potential as a high-density reactive material because of its high strength and excellent energy release performance.
Acknowledgments The authors gratefully acknowledge the help received from Mr. Y.W. Wang of the School of Materials Science and Engineering, Beijing Institute of Technology. This work was supported by the National Natural Science Foundation of China (Grant nos. 11221202, 11572049).
References [1] N.N. Thadhani, J. Appl. Phys. 76 (1994) 2129–2138. [2] Advanced Energetic Materials, The National Academies Press, vol. 32, 2004, pp. 20–23. [3] X.F. Zhang, X.N. Zhao, Chin. J. Energ. Mater. 17 (6) (2009) 731–739 (in Chinese). [4] D.E. Eakins, N.N. Thadhani, Int. Mater. Rev. 54 (2009) 181–213. [5] B.B. Aydelotte, C.H. Braithwaite, N.N. Thadhani, J. Phys.: Conf. Ser. 500 (13) (2014) 132001. [6] C.H. Braithwaite, A.L. Collins, B. Aydelotte, F. McKenzie, P.H. Chiu, in: Proceedings of the 15th Seminar on New trends in Research of Energetic Materials, 2012, pp. 93–99. [7] L.S. Bennett, F.Y. Sorrell, I.K. Simonsen, Y. Horie, K.R. Iyer, Appl. Phys. Lett. 61 (1992) 520–521. [8] D.E. Eakins, N.N. Thadhani, J. Appl. Phys. 101 (2007) 043508. [9] D.E. Eakins, N.N. Thadhani, Appl. Phys. Lett. 92 (2008) 111903. [10] B.A. Mason, L.J. Groven, S.F. Son, J. Appl. Phys. 114 (2013) 113501. [11] R.A. Austin, D.L. McDowell, D.J. Benson, Model. Simul. Mater. Sci. Eng. 22 (2014) 025018. [12] P.H. Chiu, V.F. Nesterenko, in: Proceedings of DYMAT International Conference, 2009, pp. 947–953. [13] K.L. Olney, P.H. Chiu, C.W. Lee, V.F. Nesterenko, D.J. Benson, J. Appl. Phys. 110 (11) (2011) 114908–114909. [14] V.F. Nesterenko, P.H. Chiu, C.H. Braithwaite, A. Collins, D.M. Williamson, AIP Conf. Proc. 1426 (1) (2012) 533–538. [15] K.L. Olney, D.J. Benson, V.F. Nesterenko, AIP Conf. Proc. 1426 (2012) 729–732. [16] K.L. Olney, V.F. Nesterenko, D.J. Benson, Appl. Phys. Lett. 100 (2012) 191910. [17] P.H. Chiu, K.L. Olney, C. Braithwaite, A. Jardine, A. Collins, D.J. Benson, J. Phys.: Conf. Ser. 500 (11) (2014) 112017. [18] C.T. Wei, E. Vitali, F. Jiang, S.W. Du, D.J. Benson, K.S. Vecchio, N.N. Thadhani, M. A. Meyers, Acta Mater. 60 (2012) 1418–1432. [19] B.B. Aydelotte, N.N. Thadhani, Mater. Sci. Eng. A 570 (2013) 164–171. [20] S.S. Wang, C.J. Li, X.F. Ma, J. China Ordnance 22 (2) (2001) 189–191 (in Chinese). [21] W. Chen, W.T. Zhao, J. Wang, Ordnance Mater. Sci. Eng. 32 (2) (2009) 108–110 (in Chinese). [22] X.F. Zhang, A.S. Shi, L. Qiao, J. Zhang, Y.G. Zhang, Z.W. Guan, J. Appl. Phys. 113 (2013) 083508. [23] P.G. Luo, Z.C. Wang, C.L. Jiang, L. Mao, Q. Li, Mater. Des. 84 (2015) 72–78. [24] A. Jennrich, C. Delaney, M. Clemenson, H. Krier, N. Glumac, Spring Technical Meeting of the Central States Section of the Combustion Institute, 2012, pp. 22–24. [25] A. Coverdill, C. Delaney, A. Jennrich, H. Krier, N. Glumac, J. Energ. Mater. 32 (2014) 135–145. [26] T.D. Cooper, Rockwell International, RHO-RE-ST-31P, 1984. [27] DOE Handbook, Department of Energy, DOE-HDBK, 1994, pp. 1081–1094. [28] J. Zhang, X.F. Zhang, B.Y. Fan, Ordnance Mater. Sci. Eng. 36 (1) (2013) 4–6 (in Chinese). [29] J.T. Norton, R.K. Lewis, Adv. Met. Res. Corp. (1963), N63-18389. [30] D.J. Frew, M.J. Forrestal, W. Chen, Exp. Mech. 42 (2002) 93–106. [31] D.J. Frew, M.J. Forrestal, W. Chen, Exp. Mech. 41 (2001) 40–46. [32] A.K. Kuriakose, J.L. Margrave, J. Electrochem. Soc. 111 (7) (1964) 827–831. [33] R.Z. Valiew, A.K. Mukheijee, Scr. Mater. 44 (8–9) (2001) 1747–1750.
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure and mechanical properties of W-Zr reactive materials Huilan Ren, Xiaojun Liu, Jianguo Ning n State Key Laboratory of Explosion Science and Technology, Beijing Institute of Technology, Beijing 100081, PR China
art ic l e i nf o
a b s t r a c t
Article history: Received 20 November 2015 Received in revised form 13 January 2016 Accepted 4 February 2016
Three different batches of tungsten-zirconium (W-Zr) reactive material were prepared by hot-pressing elemental powder mixtures. The first sample had a Zr:W mass ratio of 66:34. The second used zirconium hydride (ZrH2) with a (ZrH2):W mass ratio of 66:34. The third had a Zr:W mass ratio 43:57. These batches were numbered #1 to #3. To investigate the mechanical properties of the W-Zr reactive materials, samples were subjected to different strain rates using a materials testing machine and a modified split Hopkinson pressure bar (SHPB). The quasi-static compressive strengths of the three sample batches all exceeded 1022 MPa, with batch #1 sample having the highest value at about 1880 MPa. This could be ascribed to the sample exhibiting a more transgranular and dimpled fracture in the W2Zr intermetallic phase, as demonstrated by the microstructure of the fracture surface, observed by scanning electron microscopy (SEM) using an energy-dispersive spectrometer (EDS). To perform a constant strain rate experiment on this high-strength but brittle material, ramp loading using copper sheet was adopted in this study. All of the samples produced strong, bright flames when subjected to shock loading and exhibited a high compressive strength of approximately 1060 to 2690 MPa. High-resolution X-ray diffraction (XRD) was performed on the original samples and residues after the SHPB test, showing that the Zr and ZrC phases of the batch #1 and batch #3 samples, and the ZrC0.32H1.2 phase of the batch #2 sample are the active components in the reaction with air. Some small balls of ZrO2 reaction product were found not to exhibit any crystalline tungsten oxide on the residue surface. These results suggest that the batch #1 reactive material has huge potential to take the place of inert steel because of its high strength and high energy level, as well as having a density close to that of steel. & 2016 Elsevier B.V. All rights reserved.
Keywords: W-Zr Reactive materials Mechanical properties Microstructure Fracture
1. Introduction Reactive materials are a special class of energetic materials with a high energy density, which react with themselves or combust violently with air and rapidly generate high pressures and temperatures when a certain reaction initiation energy level is achieved, leading to additional damage. The reactive materials can be adapted to a variety of applications such as target damage through the addition of structural reactives, propellant/explosive additives, and manufacturing. Typical reactive materials usually consist of two or more solid-state reactants, unlike explosives that are highly sensitive to ignition, that together form a thermo-chemical mixture. These materials mainly consist of thermites, intermetallics, metal/fluorine systems, metastable intermolecular composites (MIC), nano-laminates, and metal hydrides [1–3]. Among these, the intermetallic-forming reactive materials are not only energy reactants, but can also act as structural components due to their good mechanical properties. Due to their unique n
Corresponding author. E-mail address: [email protected] (J. Ning).
http://dx.doi.org/10.1016/j.msea.2016.02.009 0921-5093/& 2016 Elsevier B.V. All rights reserved.
characteristics, the applications of intermetallics have recently diversified. They have been applied to the likes of missiles, given that their strength allows them to be used as a portion of the missile casing. The energy stored within the material will be released when subjected to the loading of impact and thus result in an increase in the explosive power of the missile, relative to that of a design using a traditional steel casing [4–6]. Intermetallics are an important subclass of reactive materials that have been studied by numerous researchers over the last several decades. They include Ni-Al and Al-W powder systems, and their study has included their initiation and characterization. The ultrafast reactions of Ni-Al powders when subjected to shock loading were first observed by Bennett et al. [7]. They found that the particle size and morphology had a significant influence on the shock ignition behavior of the Ni-Al powder mixtures in plate impact tests. This was subsequently proven by Eakins and Thadhani [8,9]. A modified Asay shear impact experiment showed that the mechanical dry milling time of the Ni-Al composites could significantly affect the microstructure and the mechanical impact ignition threshold [10]. The results of numerical simulations showed that the configuration of the Ni-Al powders played an important role in the coupled thermal–mechanical behavior under
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shock wave loading [11]. The dynamic behavior and fracturing of Al-W granular composites have been investigated and the results revealed that the morphology of the W particles had a significant effect on the dynamic behavior [12–14]. Explosively driven fragmentation experiments were performed on Al-W rings. Furthermore, the mesoscale fragmentation mechanisms that could give rise to enhanced pulverization were studied by Olney et al. [15– 17]. The failure mechanisms of axial splitting and shear failure were observed in the continuous phase of different explosively consolidated metal-aluminum powder mixtures in response to an impact [18,19]. The research results indicated that the W-Zr fragment had a capacity in excess of 93 W when initiating a charge and fuel tank, resulting from the large scattering angle and multifire point ignition. The initiation by W-Zr fragment has mainly a mechanics, while the other case has mainly a shock wave [20,21]. Impact-initiated experiments with a quasi-sealed test chamber were conducted with a W-Zr alloy and the results showed that the impact velocity could seriously impact the reaction behavior [22,23]. The W particles in W-Zr mechanical alloys that were normally regarded as being inert were proven to be involved in the reaction and released large amounts of energy under explosive tests [24,25]. Clouds of Zr dust with an average particle size of about 3 mm were susceptible to spontaneous ignition in air at room temperature. The affinity of Zr with oxygen was great, and it could also ignite in a carbon dioxide atmosphere at approximately 621 °C and in nitrogen at approximately 788 °C [26,27]. The dynamic compression mechanical properties of a W-Zr alloy under different dynamic loads were investigated and the results indicated that Zr flames reach very high temperatures, while W-Zr composites still retain their desirable structural properties before ultimately reacting [22,28]. In this study, to further improve the energy level and mechanical properties of reactive materials, hot-pressing techniques were adopted to prepare three different batches of W-Zr reactive material, the density of which was the same as steel, which could be applied to reactive material fragmentation weapons with the same kinetic energy. We investigated the differences between the microstructures and mechanical properties of the three batches of W-Zr reactive materials, when subjected to a range of loading rates.
2. Experimental setup The W-Zr reactive material was prepared from an elemental W powder (purity 99.9%, average particle size 3 mm), Zr (purity 99%, average particle size 37 mm) and ZrH2 (purity 99%, average particle size 3 mm) using hot-pressing. The powders were mixed in three different ways. The batch #1 and #2 samples had identical mass ratios with 34% of the W component. The batch #1 and #3 samples used the same Zr powder but at different mass ratios. The compositions and proportions are listed in Table 1.The preparation of the W-Zr reactive materials involved mixing and sintering. The process is shown in Fig. 1. During the sample preparation, the powders were mixed with absolute ethyl alcohol in a low-speed V-shaped blender for 6 h and then dried at 60 °C in an electrically heated forced-air drum-type
Tungsten Powder
Mixing
Zirconium/Zirconium Hydride Powder
Dry
Mixed Powder Ar Furnace Cooling W-Zr
Sintered
Fig. 1. W-Zr preparation process using hot-pressing.
Fig. 2. Typical hot-pressing curves for W-Zr.
oven for 24 h. Then, the dried mixtures were screened through a 200-mesh sieve and placed in a graphite die (Φ65 mm). Fig. 2 shows the temperature and pressure history in the sintering cycle, which was as follows: The pressure was increased to 20 MPa over 2 h and then held constant for 3.5 h. Thereafter, the pressure was reduced to 0 MPa at a rate of 60 MPa/h. The oven temperature was then increased to 1000 °C over 1 h, then slowly raised to 1500 °C at a rate of 500 °C/h and then held at 1500 °C for 3 h. After that, the temperature was decreased to 200 °C at a rate of 600 °C/h and then allowed to naturally cool to room temperature in the furnace. The entire sintering process was conducted in an argon atmosphere. For this study, each batch of the W-Zr reactive material was formed into two different sizes of sample, specifically, Φ10 mm 10 mm and Φ5 mm 5 mm, as shown in Fig. 3. Quasi-static compression tests were performed using a materials testing machine (Kexin WDW-300, China) for the samples of the three batches at room temperature, and dynamic tests were conducted using a split Hopkinson pressure bar for the W-Zr reactive material samples measuring Φ5 mm 5 mm. The apparatus consisted of a 200-mm striker bar, a 600-mm incident bar, and a 600-mm transmitting bar. All three bars were made of maraging steel with a diameter of 14.5 mm. Strain gauges were bonded to the midpoints of the incident bar and transmitted bar to record the stress waves propagating along the bars. Two high-strength
Table 1 Composition and proportion. Batch
Zr (37 mm)
ZrH2 (3 mm)
W (3 mm)
#1 #2 #3
66 0 43
0 66 0
34 34 57
1500oC Pressure
Fig. 3. Initial samples of W-Zr reactive material.
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Fig. 4. Modified SHPB apparatus.
steel spacers with the same generalized wave impedance as the bar were added between the specimen and the incident and transmission bars to protect the end surfaces against damage. The layout of the apparatus is shown in Fig. 4.
3. Results and discussion 3.1. Initial microstructures The microstructures of the samples were characterized by SEM with EDS (Hitachi S-4800, Japan). The obtained images are shown in Fig. 5. X-ray diffraction (Bruker D8 Advance, Germany) analyses were performed to identify the phases in the reactive materials. The results are shown in Fig. 6. The densities and theoretical maximum specific density (%TMD) of the three batches are listed in Table 2. As can be seen in Fig. 5, there are significant differences between the microstructures of the three batches, which illustrates the fact that the components are also different. Fig. 6 shows that the samples mainly consist of ZrC, W2Zr, and Zr phases in batch #1, ZrC0.32H1.2, W, and Zr phases in batch #2, and W, W2Zr, and ZrC phases in batch #3. The micrograph of the batch #1 sample shows the presence of only two phases (Fig. 5(a)). The dark regions are the Zr and ZrC phases, while the white regions are the W-Zr intermetallic component, which was detected by EDS (Fig. 7). Due to the high temperature, the graphite paper used to ensure the easy release of the mold was involved in the reaction to form a ZrC phase, which can be seen in the Zr-C phase diagram (Fig. 8(a)). The phase diagram for the W-Zr system indicates that, in that region with more than 20 wt% Zr and in which the temperature reaches between 1400 and 2210 °C, there is one intermetallic phase, that is, W2Zr (Fig. 8(b)). Therefore, the composition of the bright phase in batch #1 corresponds to W2Zr. The intermetallic phase is distributed evenly in the ZrC and Zr phases, although with a variety of shapes and sizes. The density of the batch #1 sample is over 99% TMD, which is close to the fully dense material. In the batch #2 sample, many of the irregular W particles were found to be surrounded by a continuous gray matrix, as shown in Fig. 5(b). The
Fig. 6. XRD patterns of samples: (a) #1, (b) #2, (c) #3. Table 2 Density and %TMD of W-Zr reactive Materials of three batches. Batch
Density (g/cm3)
% TMD
#1 #2 #3
8.34 8.01 9.15
99.2% 95.3% 87.5%
ZrH2 can remain stable after long-term aging at room temperature, since it quickly reacts with the oxygen in the air and forms a nanometer-thin layer of oxide that prevents further oxygen diffusion into the material. Nevertheless, this protective effect is lost when the temperature exceeds 400 °C and the change in the composition due to the oxidation can usually be neglected. Norton proposed that ZrC can be produced by heating a mixture of ZrH2 and graphite to 1800 to 2200 °C, and pointed out that long-term heating has a significant effect on the decomposition of the ZrH2
Fig. 5. SEM micrographs of samples before test: (a) #1, (b) #2, (c) #3.
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Table 3 Quasi-static compressive results for three batches of W-Zr reactive Materials.
Fig. 7. EDS result for white region of batch #1 sample.
hydrides [29]. In this experiment, however, the ZrH2 and W powders were sintered in a graphite die in an Ar gas atmosphere rather than in a vacuum. Because of the lower temperature (1500 °C) and the decomposition of a small amount of carbon, no ZrC was formed with the batch #2 sample. The gray phase ultimately proved to be ZrC0.32H1.2 and Zr that transformed from ZrH2 particles in the sintering process. Given the high level of activity of the small particles, the remaining Zr was easily agglomerated. The density of the batch #2 sample was 95.3% TMD, slightly less than that of batch #1. Relative to batch #1, batch #3 exhibited a smaller white region with many evenly distributed holes (Fig. 5(c)). The density of the sample is thus highly dependent on the sintering temperature. A sintering temperature of 1500 °C is too low for the batch #3 sample with a higher volume fraction of W. As a result, the grain growth occurs slowly and the crystal growth is incomplete, which leads to there being a large number of pores and thus the density is lower (87.5% TMD). 3.2. Fracture observation The quasi-static test results indicate that all three batches of the W-Zr reactive materials are a typical brittle material with a high level of strength that exhibits a good linear elastic stressstrain curve under quasi-static loading. The compressive strength
Batch
Strain (%)
Stress (MPa)
#1 #2 #3
1.0 0.7 0.4
1880 1200 1022
of the batch #1 sample is higher than that of the other two batches. The details are listed in Table 3. Scanning electron microscopy was used to investigate the failure mechanisms and the typical fracture surface morphologies of the samples after different levels of plastic deformation. The resulting images are shown in Fig. 9. A zigzag crack appeared in the fracture surface of the batch #1 sample, as shown in Fig. 9(a). It can be seen that the fracture surface consists of two different zones. The white circle indicates where the Zr phase exhibits a transgranular fracture. From the top view of the fracture (Fig. 9(b)), the detailed fracture morphology exhibits a few dimples with features of plastic fracture in region A and a predominantly cleavage river pattern in region B. Fig. 9(c) shows that the W2Zr intermetallic presents an aggregated and irregular polyhedron appearance with particle sizes ranging from 1 mm to 3 mm. An enlarged view of the crack propagation path in the W2Zr phase is shown in Fig. 9(d). As indicated by the arrows, cracks appeared in the crystal grains and along the grain boundaries, implying that these interactions absorb the energy of the crack propagation during the fracture process, and that the fracture mode is a mix of transgranular and intergranular. These results imply that transgranular fractures occurred in the Zr and ZrC phase, while a mixture of transgranular, intergranular, and dimple fractures occurred in the W2Zr phase, thus leading to a toughening of the reactive material and the attainment of the highest compressive strength. The microstructure of the batch #2 sample surface is shown in Fig. 9(e). Here, ZrC0.32H1.2 phase transgranular fractures can be easily observed. However, the microstructure of the shear fracture was very different from that of the tensile fracture, and a large number of ZrC0.32H1.2 fragments, with the appearance of wood chips with many burrs, can be observed on the surface, as shown in Fig. 9(f). Fig. 9(g) shows transgranular fractures in the ZrC0.32H1.2 phase and many spherical-shaped W particles with sizes of 1–3 mm in the fracture surface, as indicated by the circles. It is evident that the integrated W particles were separated from the matrix as a result of the intergranular fractures during the test. After the sintering and the reaction that generated the ZrC0.32H1.2, there was some Zr residue and transgranular fractures can be observed in the tensile fracture region (see Fig. 9(h)). The experiments showed
Fig. 8. Phase diagram: (a) Zr-C system, (b)W-Zr system.
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Fig. 9. SEM images of fracture surfaces after quasi-static tests: (a)-(d) batch #1, (e)-(h) batch #2, (i)-(l) batch #3.
that transgranular fractures with a small intergranular fraction were a characteristic of the batch #2 reactive material. Fig. 9(i) shows that the batch #3 sample had a very loose structure, and the cracks progress along the grain boundaries by connecting with other microcracks or across the grains as intergranular fractures. Some Zr particles exhibited a serious deformation in the shear fracture region, which is clearly demonstrated in Fig. 9(j). For clarification, a magnification of Fig. 9(i) is shown in Fig. 9(k). The fracture surface exhibited intergranular cracks and some W2Zr can be observed. Relative to batch #2 (Fig. 9(f)), the fracture surface of the batch #3 samples exhibits fewer cleavage planes, but more holes, as shown in Fig. 9(l).
3.3. Dynamic behavior In a dynamic test to determine the material properties, the sample is required to deform nearly uniformly at a constant strain rate under dynamically equal stress. For brittle materials that are loaded dynamically with the SHPB apparatus, ramp loading is typically employed to facilitate a constant strain rate experiment [30]. The typical original waveforms of the SHPB test with a pulse shaper (copper sheet) are shown in Fig. 10(a). The pulse shaping has an obvious effect on the elimination of the high-frequency signals and the delay in the incident pulse [30]. The incident wave
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Fig. 10. Results of pulse-shaped SHPB experiment with batch #1: (a) Strain-time signals, (b) Interface stresses.
becomes less steep and gives rise to ramp loading. For the batch #1 sample, tested with a gas chamber pressure of 0.1 MPa, for example, the rise time (from point A) of the incident wave was about 37 μs. For a time between 201 μs (point C) and 210 μs, the reflected wave signal was approximately equal to a constant value. Thereafter (point D), the reflected signal presented a steep rise, indicating that the sample has already been damaged or has failed. The strain signal of the sample gradually changed from a gentle slope to a steep slope between 138 μs (point B) and 150 μs. Its value changes with time, almost as a straight line, up until 162 μs, which shows that the sample deformed at a nearly constant strain rate during this period. The elastic wave velocity of the sample is about 4800 m/s and the characteristic time t¼ 1 μs (t is the duration of the stress wave required to propagate from one side in the sample to another side). Based on the literature, the rise time of the incident wave should be greater than 8t, which is a prerequisite condition for achieving the dynamic stress equilibrium in the brittle sample [31]. In this work, the stress σ1 at the incident bar/sample interface and the stress σ 2 at transmission bar/sample interface as obtained from a pulse-shaped SHPB experiment were as shown in Fig. 10(b). It can be seen that, after 10 μs, these interface stresses are in reasonably good agreement. This suggests that the samples were nearly in a state of dynamic stress equilibrium over most of the duration of the test. The sample deforms at a nearly constant strain rate under dynamically equilibrated stresses, which indicates that the data obtained from the pulse-shaped SHPB experiment was valid, and that the stresses and strains in the
samples could be calculated as follows:
⎧ 2C εR (t ) ⎪ ε ̇ (t ) = L0 ⎪ t ⎪ 2C ⎨ ε (t ) = [εR (t )] dt L0 0 ⎪ ⎪ A ⎪ σ (t ) = − Eε T (t ) A0 ⎩
∫
(1)
In Eq. (1), E, A, and C are the Young's modulus, cross-sectional area, and elastic bar wave velocity of the bar, respectively, and A0 and L0 are the cross-sectional area and the original length of the specimen, respectively. The reactive materials' maximum stress-strain relationships under quasi-static and high strain rates ranged from 200 to 1200 s 1. These are shown in Fig. 11(a). They accurately describe the compressive mechanical behaviors of the W-Zr reactive material. The quasi-static and dynamic stresses are plotted as a logarithmic function of the strain rate in Fig. 11(b). The results show that the compressive stress of all three batches of reactive materials, as determined by the dynamic tests, was higher than in the quasi-static tests, due to in-situ densification under dynamic conditions exceeding the fracture kinetics. The batch #1 samples had the highest compressive strength, and a similar behavior was also observed in the quasi-static tests. The best performance was obtained with the batch #1 samples, which can be explained by the fact that the intermetallic phases of the W2Zr provided a higher interface strength. Strain hardening
Fig. 11. Compression results for three batches under different strain rates: (a) Maximum stress vs. failure strain, (b) Maximum stress vs. strain rate.
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Fig. 12. Violent reaction of batch #1 sample when subjected to strong impact.
behavior under dynamic loading was observed, with the value ranging from 0.01 to 0.015. As shown in Fig. 11(a) (black squares), the stress-stain relationship followed a straight line. In the same way as with the trends with the batch #1 reactive material, the compressive strength of the batch #2 samples was higher at the high strain rate. The batch #2 samples had the same W mass ratio as batch #1, but were more brittle. However, the batch #3 samples had the weakest compressive strength and strain, which was attributed to the structure having many holes. The compressive strength at a strain rate of 470 s 1 differed slightly from that of the static condition (1060 MPa and 1022 MPa, respectively). The failure strain was 0.012 at a strain rate of 1200 s 1, which was more than three times that at a strain rate of 470 s 1. The samples failed, reacted, and rapidly released large amounts of chemical energy in the form of heat, light, and pressure when subjected to large impact loads. The reactions are shown in Fig. 12. This shows that the material has a very high energy content, with batch #1 exhibiting the greatest content, followed by batch #2, and then batch #3. This was determined from the size of the generated flame. According to the results of the XRD spectrum observations, the phase compositions of many of the residues are the same as those of the original samples, as is evident in Figs. 6 and 13. In addition, new diffraction peaks of the ZrO2 phases can be seen in the XRD pattern, which suggests that the Zr, ZrC, and ZrC0.32H1.2 phases are the active components. After the application of a strong impact, the high temperature and surface tension gave rise to a balling effect of the reaction product. As can be seen in Fig. 14, many small spherical balls with particle sizes of 1–10 mm are present in the
Fig. 13. XRD pattern of investigated W-Zr composite after SHPB test:(a) batch #1, (b) batch #2, (c) batch #3.
Fig. 14. SEM images of fracture surfaces of batch #1 sample after SHPB test.
fracture surface morphologies of the batch #1 sample after the SHPB test, and were finally confirmed to be Zr oxide through an EDS analysis. It should be noted that the Zr has a great affinity for oxygen, especially when the powder is very fine [27]. The ignition characteristics of the Zr and ZrH2 particles in different atmospheres were investigated by Cooper [26]. These results revealed that layers of Zr dust with particle diameters of less than 45 mm could be ignited at 530 °C in a nitrogen atmosphere, but ZrH2 powder with a particle diameter of less than 4.7 mm could not be ignited in a nitrogen atmosphere, even at 850 °C. This showed that, even though the ZrH2 powder had a smaller particle size, its affinity for nitrogen was weaker than that of the Zr powder. In addition, kinetic studies of the oxidation of ZrC were performed at about 600 °C, resulting in the formation of the ZrO2 phase [32]. The batch #1 and batch #2 samples had a higher mass ratio of elemental Zr compared to the batch #3 sample, while the batch #1 sample contained much more “fuel,” such as the Zr component, than the another two batches. These results also demonstrate why the batch #1 sample exhibited the highest luminosity during the SHPB test. Upon applying an impact load to a metal or alloy with a high stacking fault energy, due to the high strain rate and large deformation, the distribution of the dislocations becomes more uniform and the dislocation density increases, resulting in the grain refining effect being better and making it easier to produce nano-structural characteristics in the material [33]. Severe plastic deformation of the Zr particles, as well as a few clusters of irregular particles with average diameters of about 100 nm, can be observed in the batch #1 sample after the SHPB test, which can clearly be confirmed from the high-magnification SEM image (Fig. 15). These nano-structures appeared primarily near the
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Fig. 15. SEM images of configuration of batch #1 residue at higher magnification and nano-structure phase formed.
interface between the Zr phase and W2Zr phase, and this nano material was ultimately proven to be W2Zr as a result of EDS identification.
4. Conclusions
1. Three different batches of W-Zr reactive materials were successfully prepared using a hot-pressing process. The density of the batch #1 and batch #2 samples was in excess of 95% TMD, while that of the batch #3 sample was only about 87.5% TMD since the sintering temperature of 1500 °C was too low with a higher volume fraction of W. 2. The quasi-static compressive strengths of the three batches were higher than 1022 MPa, with the batch #1 sample having the highest value of about 1880 MPa. The microstructures of the fracture surfaces were investigated by SEM and EDS. The results showed that the predominant fracture modes of the W-Zr reactive materials were transgranular and intergranular, while exhibiting brittleness damage, while transgranular and dimple fractures were more predominant in the W2Zr intermetallic phase, leading to batch #1 sample toughening. 3. Waveform shaping technology was used to deform the samples at a nearly constant strain rate under dynamically equilibrated stresses, since the W-Zr reactive materials feature high strength but are brittle. A strain gauge was directly attached to the sample surface to acquire more accurate data, and the dynamic compressive strength was about 1060–2690 MPa. 4. All of the samples produced a clear flame when subjected to shock loading and an XRD analysis showed that the Zr, ZrC, and ZrC0.32H1.2 phases were the active components in the reaction with air, while reaction products of ZrO2 could be found in the residue surfaces. The batch #1 reactive material offers excellent potential as a high-density reactive material because of its high strength and excellent energy release performance.
Acknowledgments The authors gratefully acknowledge the help received from Mr. Y.W. Wang of the School of Materials Science and Engineering, Beijing Institute of Technology. This work was supported by the National Natural Science Foundation of China (Grant nos. 11221202, 11572049).
References [1] N.N. Thadhani, J. Appl. Phys. 76 (1994) 2129–2138. [2] Advanced Energetic Materials, The National Academies Press, vol. 32, 2004, pp. 20–23. [3] X.F. Zhang, X.N. Zhao, Chin. J. Energ. Mater. 17 (6) (2009) 731–739 (in Chinese). [4] D.E. Eakins, N.N. Thadhani, Int. Mater. Rev. 54 (2009) 181–213. [5] B.B. Aydelotte, C.H. Braithwaite, N.N. Thadhani, J. Phys.: Conf. Ser. 500 (13) (2014) 132001. [6] C.H. Braithwaite, A.L. Collins, B. Aydelotte, F. McKenzie, P.H. Chiu, in: Proceedings of the 15th Seminar on New trends in Research of Energetic Materials, 2012, pp. 93–99. [7] L.S. Bennett, F.Y. Sorrell, I.K. Simonsen, Y. Horie, K.R. Iyer, Appl. Phys. Lett. 61 (1992) 520–521. [8] D.E. Eakins, N.N. Thadhani, J. Appl. Phys. 101 (2007) 043508. [9] D.E. Eakins, N.N. Thadhani, Appl. Phys. Lett. 92 (2008) 111903. [10] B.A. Mason, L.J. Groven, S.F. Son, J. Appl. Phys. 114 (2013) 113501. [11] R.A. Austin, D.L. McDowell, D.J. Benson, Model. Simul. Mater. Sci. Eng. 22 (2014) 025018. [12] P.H. Chiu, V.F. Nesterenko, in: Proceedings of DYMAT International Conference, 2009, pp. 947–953. [13] K.L. Olney, P.H. Chiu, C.W. Lee, V.F. Nesterenko, D.J. Benson, J. Appl. Phys. 110 (11) (2011) 114908–114909. [14] V.F. Nesterenko, P.H. Chiu, C.H. Braithwaite, A. Collins, D.M. Williamson, AIP Conf. Proc. 1426 (1) (2012) 533–538. [15] K.L. Olney, D.J. Benson, V.F. Nesterenko, AIP Conf. Proc. 1426 (2012) 729–732. [16] K.L. Olney, V.F. Nesterenko, D.J. Benson, Appl. Phys. Lett. 100 (2012) 191910. [17] P.H. Chiu, K.L. Olney, C. Braithwaite, A. Jardine, A. Collins, D.J. Benson, J. Phys.: Conf. Ser. 500 (11) (2014) 112017. [18] C.T. Wei, E. Vitali, F. Jiang, S.W. Du, D.J. Benson, K.S. Vecchio, N.N. Thadhani, M. A. Meyers, Acta Mater. 60 (2012) 1418–1432. [19] B.B. Aydelotte, N.N. Thadhani, Mater. Sci. Eng. A 570 (2013) 164–171. [20] S.S. Wang, C.J. Li, X.F. Ma, J. China Ordnance 22 (2) (2001) 189–191 (in Chinese). [21] W. Chen, W.T. Zhao, J. Wang, Ordnance Mater. Sci. Eng. 32 (2) (2009) 108–110 (in Chinese). [22] X.F. Zhang, A.S. Shi, L. Qiao, J. Zhang, Y.G. Zhang, Z.W. Guan, J. Appl. Phys. 113 (2013) 083508. [23] P.G. Luo, Z.C. Wang, C.L. Jiang, L. Mao, Q. Li, Mater. Des. 84 (2015) 72–78. [24] A. Jennrich, C. Delaney, M. Clemenson, H. Krier, N. Glumac, Spring Technical Meeting of the Central States Section of the Combustion Institute, 2012, pp. 22–24. [25] A. Coverdill, C. Delaney, A. Jennrich, H. Krier, N. Glumac, J. Energ. Mater. 32 (2014) 135–145. [26] T.D. Cooper, Rockwell International, RHO-RE-ST-31P, 1984. [27] DOE Handbook, Department of Energy, DOE-HDBK, 1994, pp. 1081–1094. [28] J. Zhang, X.F. Zhang, B.Y. Fan, Ordnance Mater. Sci. Eng. 36 (1) (2013) 4–6 (in Chinese). [29] J.T. Norton, R.K. Lewis, Adv. Met. Res. Corp. (1963), N63-18389. [30] D.J. Frew, M.J. Forrestal, W. Chen, Exp. Mech. 42 (2002) 93–106. [31] D.J. Frew, M.J. Forrestal, W. Chen, Exp. Mech. 41 (2001) 40–46. [32] A.K. Kuriakose, J.L. Margrave, J. Electrochem. Soc. 111 (7) (1964) 827–831. [33] R.Z. Valiew, A.K. Mukheijee, Scr. Mater. 44 (8–9) (2001) 1747–1750.