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Preparation of W-Cu alloy with high density and ultrafine grains by mechanical alloying and high pressure sintering
Int. Journal of Refractory Metals and Hard Materials 61 (2016) 91–97
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Preparation of W-Cu alloy with high density and ultrafine grains by mechanical alloying and high pressure sintering W.T. Qiu a, Y. Pang a, Z. Xiao a, Z. Li a,b,⁎ a b
School of Materials Science and Engineering, Central South University, Changsha 410083, China Key Laboratory of Nonferrous Metal Materials Science and Engineering, Ministry of Education, Changsha 410083, China
a r t i c l e
i n f o
Article history: Received 29 August 2015 Received in revised form 4 July 2016 Accepted 8 July 2016 Available online 19 July 2016 Keywords: W-Cu alloy Mechanical alloying High pressure sintering Densification mechanism
a b s t r a c t W-20 wt%Cu alloy, with a high sintering density of 99.2% theoretical density (%TD), ultrafine grain size of 500 nm, and hardness of 310 HV, was fabricated by mechanical alloying and high pressure sintering (HPS). The changes of morphology and structural evolution of W-Cu powders during mechanical alloying process were observed and studied, and the effects of temperature and pressure on the density of sintered compacts were studied. Results showed that nanostructured (NS) mechanical alloyed (MA) W-Cu powders were obtained after a 40 h milling at a milling speed of 400 rpm, and the W-Cu alloy had nanocrystalline structure and shape retention with nearly full densification when the HPS of NS MA W-Cu powders was carried out at 1050 °C with a pressure of 40 MPa. The densification behavior of NS MA W-Cu powders during HPS was also studied, including a nanosintering within the MA powders during heating and a global rearrangement and sintering when an external pressure was applied. © 2016 Elsevier Ltd. All rights reserved.
1. Introduction Due to the low thermal expansion coefficient of tungsten (W) and high electrical conductivity of copper (Cu), W-Cu alloys have been widely used as microwave materials, diverter plate materials in fusion reactor materials and warhead materials [1–4]. Conventionally, liquid infiltration sintering (LIS) and liquid phase sintering (LPS) are used to prepare W-Cu alloys. However, it is difficult to attain full density WCu alloys using these procedures because W-Cu system exhibits mutual insolubility or negligible solubility [5,6]. The sinterablility of W-Cu powders can be improved by adding Ni, Fe or Co as sintering activators [7], whereas the extra elements generally decrease the electrical and thermal conductivity of W-Cu alloys [8,9]. Hot-shock consolidation (HSC) was used to fabricate fine-grained W-Cu alloy under a short sintering time (b 3 min) without any sintering activator [10]. However, the operations of HSC were very complicated and the microstructural homogeneity was difficult to control by HSC. The fabrication of W-Cu alloy by sintering Cu-coated W powders was also introduced [11], but formaldehyde, as the reducing agent during coating, has some disadvantages such as being carcinogenic and irritative to the human body. Comparing to aforementioned methods, the mechanical alloying process is more attractive since it not only is able to enhance microstructural homogeneity, but also offers advantages for the preparation ⁎ Corresponding author at: School of Materials Science and Engineering, Central South University, Changsha 410083, China. E-mail address: [email protected] (Z. Li).
http://dx.doi.org/10.1016/j.ijrmhm.2016.07.013 0263-4368/© 2016 Elsevier Ltd. All rights reserved.
of nanostructured (NS) materials [12–16]. It is well known that the well-mixed NS powders can improve the sinterability, which may lead to high density of W-Cu alloy. To obtain the nearly full density W-Cu alloys, the sintering process of the NS mechanical alloyed (MA) W-Cu powders is usually carried out at a temperature above the melting point of Cu since the rearrangement of MA powders can be induced by the capillary force of liquid Cu phase. However, the nano-sized W grain obtained by mechanical alloying process could grow up to several micrometers during liquid phase sintering. To obtain W-Cu alloy with both high density and ultrafine grains, high pressure sintering (HPS) was introduced. The sintering temperatures were set near the melting point of Cu. Applying high pressure on the compact and choosing relatively low temperature during sintering offer several advantages. Firstly, the Cu phase under this condition shows unique properties of both solid phase and liquid phase. It can not only fill in gaps among MA powders but also avoid being extruded from the die under the high pressure, which can maintain the homogeneity of the compact with high density. More importantly, ultrafine grains can be obtained due to the inhibition of W grain growth under this condition. All of these advantages would greatly enhance the properties of final W-Cu alloy products. In this study, W-20 wt% Cu alloys with high density and ultra-fine W grains were obtained by mechanical alloying and HPS at a temperature of 1050 °C, near to the melting point of Cu. The changes of morphology and structural evolution of W-Cu powders during the mechanical alloying process were investigated, and the effects of temperature and pressure on the density of sintered compacts were also studied. Finally,
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(a)
(b)
5µm
5µm 100
100
(c)
*
(d)
D (v, 0.9) = 17.90 um
* D (v, 0.9) = 20.15 um
80
Volume(%)
Volume (%)
80
60
* D (v, 0.5) = 8.84 um
40
60
*
D (v, 0.5) = 10.83 um
40
20
20
* D (v, 0.1) = 5.82 um
* D (v, 0.1) = 4.60 um
0 0
10
20
0
30
40
Size (um)
50
0
10
20
30
40
50
Size (um)
Fig. 1. SEM images and size distributions of the W ((a) and (c)) and Cu powders ((b) and (d)).
the mechanism for obtaining ultra-fine W grains during both mechanical alloying and HPS process, as well as the densification mechanism of NS MA W-Cu powders during HPS were discussed. 2. Experimental procedures 2.1. Mechanical alloying process Pure W (99.99% purity) and Cu (99.99% purity) powders as shown in Fig. 1 were used. It can be seen that W particles exhibited polyhedral shape (Fig. 1(a)), while the majority of the Cu particles were near spherical (Fig. 1(b)). Both W and Cu particles showed agglomeration. Therefore, the particle sizes of W and Cu powders detected by the Mastersizer Micro particle size analyzer (Fig. 1(c) and (d)) were much larger than those shown in SEM images. Before the mechanical alloying process, W and Cu powders were reduced under hydrogen at 800 °C for 1 h and at 400 °C for 45 min, respectively. Then they were mixed for 30 min in a V-type mixing machine with the composition corresponding to W-20 wt% Cu. The mechanical alloying process was carried out in a QM-1F highenergy planetary ball mill (Nanjing, China) with four stainless steel vessels (1000 ml each). Each vial contained hardened steel balls of different sizes (20 mm, 10 mm and 6 mm in diameter) with a ball-to-powder weight ratio of 15:1. The sealed vials were evacuated and then filled with argon to avoid oxidation of the powders during the mechanical alloying process. At different stages of milling (5, 15, 40, 60, 80 and 100 h), a small batch of the milled powders was taken out of the vials for analysis and testing.
pressure of 100 ± 10 MPa to produce green parts with 50 ± 2%TD. HPS was performed at different temperatures and pressures, ranging from 900 °C to 1100 °C and 0 MPa to 50 MPa, respectively, for 2 h. The heating rate was 10 °C/min, and the vacuum was kept below 10−3 Pa. For comparison, a W-20 wt%Cu powder mixture (obtained by directly mixing W and Cu powders in a V-type mixing machine with a speed of 30 rpm for 5 h) was also sintered by HPS. 2.3. Techniques used to characterize the powders and compacts The characteristics of the MA W-Cu powders were investigated by Xray diffraction (XRD) using a Dmax-2500 diffractometer with Cu Kα radiation. The intensity vs. Bragg angle data for all XRD samples was collected at a 2θ from 35° to 120° with a 0.025° step size and under a dwell time of 5 s. The morphologies of the MA powders and sintered compacts were observed using a Sirion 200 scanning electron microscope equipped with EDAX GENESIS 60, and the chemical compositions of MA powders were analyzed using inductively coupled plasma optical emission spectrophotometer (ICP-OES). The phase transformation within the NS MA W-Cu powders was examined by differential thermal analysis using an STA 449C simultaneous thermal analyzer (NETZSCH, Germany). The densities of the sintered compacts were measured by Archimedes displacement method. Vickers hardness of the compacts was measured under a load of 20 Kp. In order to study the binding condition among W grains, specimens were taken from the sintered compacts and immersed in a 50% nitric acid solution to remove the Cu, then washed, and finally dried in hydrogen. 3. Results and discussions
2.2. High pressure sintering process 3.1. Microstructural characteristics of as-milled powders HPS process was carried out in a vacuum hot press sintering furnace (Shenyang, China), which consisted of a 30 t auto-operated hydraulic press and a two-action piston die of 30 mm bore diameter. Before HPS process, the MA W-Cu powders were compacted in a cylindrical die at a
Fig. 2 shows representative SEM images of MA W-Cu powders milled for different periods. It can be seen from Fig. 2(a) that the particle size of the powders milled for 5 h was larger than that of raw powders due to
W.T. Qiu et al. / Int. Journal of Refractory Metals and Hard Materials 61 (2016) 91–97
(a)
93
(b)
10µm
10µm
(d)
(c)
10µm
10µm
(e)
(f)
200nm
200nm
Fig. 2. SEM images of MA W-Cu powders milled for different periods (5 h (a), 15 h (b), 40 h (c) (e) and 60 h (d) (f)).
the cold welding effect. Both W and Cu grains (shown in Fig. 1) were crushed during the mechanical alloying process (Fig. 2(a)). With longer milling time, the particle size decreased and the shape of composite particles became near spherical (Fig. 2(b)). When the milling time increased to 40 h, the size of particles became nearly stable. Welding, fracturing and re-welding of the elemental powders occurred repeatedly during mechanical alloying process (Fig. 2(c) and (d)). From the SEM images with higher magnification (Fig. 2(e) and (f)), it can be
seen that the particles consisted of many nano-sized grains, and the size of the gains was decreased with the increase of the milling time. However, as shown in Table 1, considerable amount of impurities can be introduced if too long milling time is applied, in particular of Fe, a common occurrence during ball milling [13]. It has been found that high level of impurities had significant negative influence on the properties of W-Cu alloys, especially for thermal conductivity [8,9,17]. Therefore, the MA W-Cu powders milled for 40 h, with relatively low impurities level, were chosen to fabricate W-Cu alloys by HPS.
Table 1 Chemical analysis of MA W-Cu powders milled for different periods.
3.2. Densification results of HPS compacts
Milling time/hour
W
Cu
Fe
Cr
Ni
(wt%) Mn
5 15 40 60 80 100
82.13 78.42 80.03 79.93 79.98 80.03
17.81 21.48 19.79 19.71 19.51 19.14
0.033 0.063 0.118 0.263 0.336 0.574
0.019 0.025 0.048 0.084 0.153 0.231
0.004 0.006 0.007 0.009 0.012 0.018
0.001 0.002 0.004 0.006 0.007 0.009
The effects of temperature and pressure on the density and hardness of the sintered compacts were studied. Fig. 3(a) displays the densities and hardness of the compacts prepared by HPS of NS MA W-Cu powders at different temperatures under a pressure of 50 MPa. The densities are normalized to the theoretical density of W-20 wt% Cu of 15.64 g/cm3. As can be seen in Fig. 3(a), the densification of the compact was strongly influenced by the sintering temperature. From a density of 85.11%TD obtained at a sintering temperature of 900 °C, the density increased
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100
1100 oC
(b)
95 300 90 250
Hardness (HV)
Relative density (%TD)
350
relative density hardness
(a)
(c)
85 200 80
pores 150
75 950 1000 1050 o Sintering temperature ( C)
1100
350
100
Relative density (%TD)
(d)
relative density hardness
95 300 90 250
Hardness (HV)
900
85 200 80
150
75 0
20
30 Pressure (MPa)
40
50
Fig. 3. The bar graph illustrations for densities and hardness of the W-Cu sintered compacts prepared by HPS of NS MA W-Cu powders at different temperatures under a pressure of 50 MPa (a); the macroscopic photo of the die (b) and the SEM image of the W-Cu sintered compact (c) when HPS was carried out at 1100 °C under a pressure of 50 MPa; densities and hardness of the W-Cu sintered compacts prepared by HPS of NS MA W-Cu powders under different pressures at 1050 °C (d).
with sintering temperature, and reached the highest value (99.1%TD) at a sintering temperature of 1050 °C. An attempt to increase the sintering temperature into the liquid phase led to the loss of Cu melt out of the die (Fig. 3(b)), and to a high porosity within the samples (Fig. 3(c)). Fig. 3(d) displayed the densities and hardness of the compacts prepared by HPS of NS MA W-Cu powders under different pressures at
(a)
1050 °C. As shown, the pressure played an important role in the densification of compacts during HPS. The density of compact was only 86.2%TD without pressure (P = 0 MPa), and rose to 92.4%TD when a pressure of 20 MPa was applied. It reached 99.2%TD under a pressure of 40 MPa, and then remained nearly stable when the pressure was further increased.
(b)
copper pools
pores 10µm
10µm
Fig. 4. SEM images of W-Cu sintered compacts prepared from NS MA W-Cu powders (a) and standard mixed W-Cu powders (b).
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95
3.3. Properties of W-Cu alloy
100 hours CPS
The properties of the W-Cu alloy sintered at 1050 °C under a pressure of 40 MPa were investigated. Figs. 4(a) and (b) show SEM images of W-Cu sintered compacts prepared from NS MA W-Cu powders and that of the reference sample obtained from mixed W-Cu powders respectively. In Fig. 4, the light grey area is the W phase, while Cu appears as dark grey areas, and the pores appear as black areas. Compared to the compact prepared from mixed W-Cu powders, the compact prepared form NS MA W-Cu powders exhibited fewer pores and a more homogeneous microstructure consisting of ultra-fine W grains and a continuous Cu network, and some Cu rich pools outside the W particles (marked by the red arrows). Consistent with the morphologies shown in Fig. 4, the density of the compact prepared from NS MA W-Cu powder (99.2%TD) was much higher than that achieved using the mixed W-Cu powder (90%TD). After removal of Cu, the specimen prepared from mixed W-Cu powders collapsed, while the specimen prepared from NS MA W-Cu powders retained its shape. A SEM image of this sample is shown in Fig. 5. Although the sintering temperature was relatively low regarding W, sintering necks between W grains can be observed. This is likely to be the reason for the mechanical stability of the sample. Because of the low sintering temperature, the W grains coarsened slightly and had submicron size (about 500 nm). Consistent with the low porosity, the average hardness of the W-Cu sintered compacts prepared from NS MA powders reached close to 310 HV, much higher than that prepared from mixed W and Cu powders (average 260 HV).
80 hours 60 hours 40 hours 15 hours 5 hours W Cu
45
4.1. The effect of mechanical alloying process on the sintering properties of powders Fig. 6 shows the XRD patterns of MA W-Cu powders milled for different periods. With the increase of milling time, intensities of the diffraction peaks of Cu decreased rapidly and almost disappeared after 40 h milling. After milling for 100 h, only W diffraction peaks can be observed superposed on a broad background. A progressive broadening of the W peaks was observed with the increase of the milling time, which could be induced by grain refinement and lattice strain accumulation [16]. Generally, there are five factors contributing to the broadening of peaks appearing in XRD patterns: (1) crystallite size; (2) micro-strain; (3) diffraction instrument; (4) the overlapping of two sets of diffraction profiles from Kα1 and Kα2 diffraction of Cu; (5) scattering of background [18]. A standard Si sample was used to deal with the intrinsic instrument broadening effect. The effect of Cu Kα2 radiation on the peak was eliminated by Fourier series separation, and the scattering of background could be evaluated by the quantic polynomial expression.
(a)
W
60
75 o 2 ( )
W
W
90
105
W
120
Fig. 6. XRD patterns of MA W-Cu powders milled for different periods.
Then the crystallites size and lattice micro-strain can be related to the width of peaks by Williamson–Hall method [19]: β cosθ ¼
4. Discussions
Cu
W
Kλ þ 2ε sinθ d
ð1Þ
Where β is the full width at half-maximum (FWHM) of the MA powders, θ is the position of peak maximum, K is the Scherrer constant (about 0.9), λ is the radiation wavelength (about 1.5405 Å), d is the crystallite size and ε is the lattice micro-strain introduced by milling. Fig. 7 showed the variations of calculated crystallite size and microstrain of the W matrix with different milling times. The average crystallite size dropped rapidly in the first 5 h, and then dropped gradually to about 10.9 nm when the milling time increased to 40 h. After 40 h, the average crystallite sizes kept nearly constant, and the finest average crystallite size of 7.1 nm was attained at 100 h. In contrast to the variation of crystallite size, the micro-strain rose rapidly at the initial stage, and then increased gently after 40 h milling. As shown above, the MA W-Cu powders milled for 40 h exhibited a nano-crystalline microstructure and a large micro-strain, which obviously enhances the sintering properties of the powders [20]. We hold this is the major reason that the density of the sintered compact prepared from NS MA W-Cu powders is much higher than that prepared by simply mixed W-Cu powders.
(b)
10µm Fig. 5. SEM images of W skeleton after removal of Cu of the simple prepared from NS MA W-Cu powders.
500nm
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0.7
100
300
0.6
80
0.5
250
Crystallite size Micro-strain
0.4
200 0.3 150 0.2
Heat flow (W/g)
350
Micro-strain (%)
Crystallite size (nm)
96
60 40 20 0
100
1083 oC
0.1
50
-20
0.0
Exothermic -40
-0.1 0
20
40 60 Milling time (hour)
80
100
200
400
600
800
1000
Temperature (oC)
Fig. 7. Effects of milling time on crystallite size and micro-strain of the W matrix. Fig. 8. DTA curve of NS MA W-Cu powders at a heating rate of 10 °C/minute.
4.2. Densification mechanism of NS MA W-Cu powders during high pressure sintering The densification mechanism of NS W-Cu powders prepared by mechanical alloying process was different from that of the standard mixed W-Cu powders because of their different characteristic morphologies. As has been discussed in previous studies [20–28], the densification mechanism of NS MA W-Cu powders during sintering can be described as follows: the Cu-pools and Cu-layers are formed by the growth of Wgrains inside MA powders at the temperature near the melting point of Cu, referred to as “nanosintering phenomenon”, and the Cu phase diffuses out to the surface of individual MA powder particles and surrounds the exterior of MA powders, leading to a bonding between the particles. Further increase in density requires the rearrangement of particles. To induce this, increasing the temperature beyond the melting point of Cu is the most effective method, as the MA powder particles can be rearranged by the capillary force of liquid of the preformed Cupools and Cu-layers. It was reported that about 95%TD can be achieved by the particle rearrangement during liquid phase-sintering [21]. And Heady and Cahn [29] proposed that the maximum interparticle force due to capillary action was roughly equivalent to an external pressure: pffiffiffi P ¼ 2 2π ðγlv =RÞ cosθ
ð2Þ
Where γlv denotes the surface free energy of the liquid-vapor, R is radius of the particles, θ is the contact angle of liquid-solid, and P represents the equivalent external pressure. The higher the temperature, the smaller the θ. Therefore, full densification of standard mixed W-Cu powders was attained at temperatures above 1350 °C because a complete wetting of W-Cu system was then obtained, whereas full densification of NS MA W-Cu powders was already attained at about 1200 °C [21,22]. Fig. 8 showed the differential thermal analysis (DTA) image of NS MA W-Cu powders milled for 40 h. It can be seen that the melting point of Cu in the powders was about 1083 °C. The HPS process of NS MA powders was carried out at 1050 °C, below the melting point of Cu. Hence, the rearrangement of MA powder particles by capillary force could not be fully obtained. However, when an external pressure of 20 MPa was applied, the density of compact rose from 86.2%TD to 92.4%TD, and nearly full density (99.2%TD) could be obtained when the pressure was increased to 40 MPa, indicating that the rearrangement of MA powder particles can be driven by the external high pressure. The role of high pressure played in the densification of MA powders during sintering is equally important as the capillary force in liquid phase sintering. As a result, nearly full densification of MA powders can be achieved already at relatively low sintering temperature
under high pressure on the compact. Meanwhile, the temperature for HPS should be not too low. When the sintering temperature of HPS was much lower than the melting point of Cu, i.e., 950 °C, the density was still lower than 90%TD even at a pressure of 50 MPa (as shown in Fig. 4(a)). This we attribute to the fact that the rearrangement of MA powder particles is hindered in the absence of Cu-pools and Cu-layers which are usually formed at 1000 °C [21]. When the sintering temperature increased to 1050 °C, not only the Cu pools and layers were formed, but also the Cu phase became viscous and attained a certain extent of mobility at the temperature near the melting point of Cu [24]. Therefore, the rearrangement of MA powder particles was facilitated, and the gaps between MA powders could be filled with Cu phase when a pressure of 40 MPa was applied, resulting in nearly full density of the sintered compact. Because of the relatively low sintering temperature, the average size of W grains was about 500 nm, which was smaller than that of W grains sintered at the temperature of liquid phase (nearly or more than 1 μm) [21–23]. According to the study by Ryu and Kim, the bonding of MA powders sintered below the melting point of Cu was formed by the diffusion of Cu phase [20]. However, the bonding of MA powders during HPS was accomplished also by the sintering of W grains, as described in previous section. There are three main concerns on sintering of W– Cu nanocomposites: Complete densification, maintenance of nanocrystalline structure and shape retention [30,31]. Compared with liquid phase sintering, HPS is more conducive to achieve nanocrystalline structure and shape retention with nearly full densification. 5. Conclusions W-20 wt% Cu alloy with a high density (99.2%TD) and ultra-fine W grains (500 nm) can be successfully prepared by the HPS at 1050 °C under a pressure of 40 MPa. The changes of morphology and structural evolution of W-Cu powders during mechanical alloying process were observed and studied, and the effects of temperatures and pressure on the density of sintered compacts made from NS MA W-Cu powders were investigated. Based on these studies, the following conclusions can be made: 1. Nanosized MA W-Cu powders can be obtained after 40 h milling at a milling speed of 400 rpm. The results of XRD analysis indicate that these MA powders are nanocrystalline with a grain size around 10 nm. Extending the milling time did not reduce the W grain size significantly, but increased the impurity content, in particular of Fe, in the MA W-Cu powders. 2. The optimum temperature and pressure for HPS of NS MA W-Cu powders were 1050 °C and 40 MPa, respectively. At this temperature,
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ultra-fine W grains can be obtained in the final samples. The densification of MA powders could be strongly promoted due to sintering under pressure. The density increased with pressure, going into a limit at about 99%TD. 3. The densification of NS MA W-Cu powders during HPS at a temperature near the melting point of Cu is attributed to a nanosintering process accompanied by the rearrangement of MA powder particles under external high pressure. Acknowledgements The authors gratefully acknowledge financial support from the Natural Science Foundation of China (51271203), Self-dependent Innovation Project for Post-graduate Students of Central South University (2012zzts006), and grants from the project of innovation-driven plan in Central South University.
[12] [13]
[14]
[15]
[16] [17] [18]
[19] [20]
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Int. Journal of Refractory Metals and Hard Materials journal homepage: www.elsevier.com/locate/IJRMHM
Preparation of W-Cu alloy with high density and ultrafine grains by mechanical alloying and high pressure sintering W.T. Qiu a, Y. Pang a, Z. Xiao a, Z. Li a,b,⁎ a b
School of Materials Science and Engineering, Central South University, Changsha 410083, China Key Laboratory of Nonferrous Metal Materials Science and Engineering, Ministry of Education, Changsha 410083, China
a r t i c l e
i n f o
Article history: Received 29 August 2015 Received in revised form 4 July 2016 Accepted 8 July 2016 Available online 19 July 2016 Keywords: W-Cu alloy Mechanical alloying High pressure sintering Densification mechanism
a b s t r a c t W-20 wt%Cu alloy, with a high sintering density of 99.2% theoretical density (%TD), ultrafine grain size of 500 nm, and hardness of 310 HV, was fabricated by mechanical alloying and high pressure sintering (HPS). The changes of morphology and structural evolution of W-Cu powders during mechanical alloying process were observed and studied, and the effects of temperature and pressure on the density of sintered compacts were studied. Results showed that nanostructured (NS) mechanical alloyed (MA) W-Cu powders were obtained after a 40 h milling at a milling speed of 400 rpm, and the W-Cu alloy had nanocrystalline structure and shape retention with nearly full densification when the HPS of NS MA W-Cu powders was carried out at 1050 °C with a pressure of 40 MPa. The densification behavior of NS MA W-Cu powders during HPS was also studied, including a nanosintering within the MA powders during heating and a global rearrangement and sintering when an external pressure was applied. © 2016 Elsevier Ltd. All rights reserved.
1. Introduction Due to the low thermal expansion coefficient of tungsten (W) and high electrical conductivity of copper (Cu), W-Cu alloys have been widely used as microwave materials, diverter plate materials in fusion reactor materials and warhead materials [1–4]. Conventionally, liquid infiltration sintering (LIS) and liquid phase sintering (LPS) are used to prepare W-Cu alloys. However, it is difficult to attain full density WCu alloys using these procedures because W-Cu system exhibits mutual insolubility or negligible solubility [5,6]. The sinterablility of W-Cu powders can be improved by adding Ni, Fe or Co as sintering activators [7], whereas the extra elements generally decrease the electrical and thermal conductivity of W-Cu alloys [8,9]. Hot-shock consolidation (HSC) was used to fabricate fine-grained W-Cu alloy under a short sintering time (b 3 min) without any sintering activator [10]. However, the operations of HSC were very complicated and the microstructural homogeneity was difficult to control by HSC. The fabrication of W-Cu alloy by sintering Cu-coated W powders was also introduced [11], but formaldehyde, as the reducing agent during coating, has some disadvantages such as being carcinogenic and irritative to the human body. Comparing to aforementioned methods, the mechanical alloying process is more attractive since it not only is able to enhance microstructural homogeneity, but also offers advantages for the preparation ⁎ Corresponding author at: School of Materials Science and Engineering, Central South University, Changsha 410083, China. E-mail address: [email protected] (Z. Li).
http://dx.doi.org/10.1016/j.ijrmhm.2016.07.013 0263-4368/© 2016 Elsevier Ltd. All rights reserved.
of nanostructured (NS) materials [12–16]. It is well known that the well-mixed NS powders can improve the sinterability, which may lead to high density of W-Cu alloy. To obtain the nearly full density W-Cu alloys, the sintering process of the NS mechanical alloyed (MA) W-Cu powders is usually carried out at a temperature above the melting point of Cu since the rearrangement of MA powders can be induced by the capillary force of liquid Cu phase. However, the nano-sized W grain obtained by mechanical alloying process could grow up to several micrometers during liquid phase sintering. To obtain W-Cu alloy with both high density and ultrafine grains, high pressure sintering (HPS) was introduced. The sintering temperatures were set near the melting point of Cu. Applying high pressure on the compact and choosing relatively low temperature during sintering offer several advantages. Firstly, the Cu phase under this condition shows unique properties of both solid phase and liquid phase. It can not only fill in gaps among MA powders but also avoid being extruded from the die under the high pressure, which can maintain the homogeneity of the compact with high density. More importantly, ultrafine grains can be obtained due to the inhibition of W grain growth under this condition. All of these advantages would greatly enhance the properties of final W-Cu alloy products. In this study, W-20 wt% Cu alloys with high density and ultra-fine W grains were obtained by mechanical alloying and HPS at a temperature of 1050 °C, near to the melting point of Cu. The changes of morphology and structural evolution of W-Cu powders during the mechanical alloying process were investigated, and the effects of temperature and pressure on the density of sintered compacts were also studied. Finally,
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(a)
(b)
5µm
5µm 100
100
(c)
*
(d)
D (v, 0.9) = 17.90 um
* D (v, 0.9) = 20.15 um
80
Volume(%)
Volume (%)
80
60
* D (v, 0.5) = 8.84 um
40
60
*
D (v, 0.5) = 10.83 um
40
20
20
* D (v, 0.1) = 5.82 um
* D (v, 0.1) = 4.60 um
0 0
10
20
0
30
40
Size (um)
50
0
10
20
30
40
50
Size (um)
Fig. 1. SEM images and size distributions of the W ((a) and (c)) and Cu powders ((b) and (d)).
the mechanism for obtaining ultra-fine W grains during both mechanical alloying and HPS process, as well as the densification mechanism of NS MA W-Cu powders during HPS were discussed. 2. Experimental procedures 2.1. Mechanical alloying process Pure W (99.99% purity) and Cu (99.99% purity) powders as shown in Fig. 1 were used. It can be seen that W particles exhibited polyhedral shape (Fig. 1(a)), while the majority of the Cu particles were near spherical (Fig. 1(b)). Both W and Cu particles showed agglomeration. Therefore, the particle sizes of W and Cu powders detected by the Mastersizer Micro particle size analyzer (Fig. 1(c) and (d)) were much larger than those shown in SEM images. Before the mechanical alloying process, W and Cu powders were reduced under hydrogen at 800 °C for 1 h and at 400 °C for 45 min, respectively. Then they were mixed for 30 min in a V-type mixing machine with the composition corresponding to W-20 wt% Cu. The mechanical alloying process was carried out in a QM-1F highenergy planetary ball mill (Nanjing, China) with four stainless steel vessels (1000 ml each). Each vial contained hardened steel balls of different sizes (20 mm, 10 mm and 6 mm in diameter) with a ball-to-powder weight ratio of 15:1. The sealed vials were evacuated and then filled with argon to avoid oxidation of the powders during the mechanical alloying process. At different stages of milling (5, 15, 40, 60, 80 and 100 h), a small batch of the milled powders was taken out of the vials for analysis and testing.
pressure of 100 ± 10 MPa to produce green parts with 50 ± 2%TD. HPS was performed at different temperatures and pressures, ranging from 900 °C to 1100 °C and 0 MPa to 50 MPa, respectively, for 2 h. The heating rate was 10 °C/min, and the vacuum was kept below 10−3 Pa. For comparison, a W-20 wt%Cu powder mixture (obtained by directly mixing W and Cu powders in a V-type mixing machine with a speed of 30 rpm for 5 h) was also sintered by HPS. 2.3. Techniques used to characterize the powders and compacts The characteristics of the MA W-Cu powders were investigated by Xray diffraction (XRD) using a Dmax-2500 diffractometer with Cu Kα radiation. The intensity vs. Bragg angle data for all XRD samples was collected at a 2θ from 35° to 120° with a 0.025° step size and under a dwell time of 5 s. The morphologies of the MA powders and sintered compacts were observed using a Sirion 200 scanning electron microscope equipped with EDAX GENESIS 60, and the chemical compositions of MA powders were analyzed using inductively coupled plasma optical emission spectrophotometer (ICP-OES). The phase transformation within the NS MA W-Cu powders was examined by differential thermal analysis using an STA 449C simultaneous thermal analyzer (NETZSCH, Germany). The densities of the sintered compacts were measured by Archimedes displacement method. Vickers hardness of the compacts was measured under a load of 20 Kp. In order to study the binding condition among W grains, specimens were taken from the sintered compacts and immersed in a 50% nitric acid solution to remove the Cu, then washed, and finally dried in hydrogen. 3. Results and discussions
2.2. High pressure sintering process 3.1. Microstructural characteristics of as-milled powders HPS process was carried out in a vacuum hot press sintering furnace (Shenyang, China), which consisted of a 30 t auto-operated hydraulic press and a two-action piston die of 30 mm bore diameter. Before HPS process, the MA W-Cu powders were compacted in a cylindrical die at a
Fig. 2 shows representative SEM images of MA W-Cu powders milled for different periods. It can be seen from Fig. 2(a) that the particle size of the powders milled for 5 h was larger than that of raw powders due to
W.T. Qiu et al. / Int. Journal of Refractory Metals and Hard Materials 61 (2016) 91–97
(a)
93
(b)
10µm
10µm
(d)
(c)
10µm
10µm
(e)
(f)
200nm
200nm
Fig. 2. SEM images of MA W-Cu powders milled for different periods (5 h (a), 15 h (b), 40 h (c) (e) and 60 h (d) (f)).
the cold welding effect. Both W and Cu grains (shown in Fig. 1) were crushed during the mechanical alloying process (Fig. 2(a)). With longer milling time, the particle size decreased and the shape of composite particles became near spherical (Fig. 2(b)). When the milling time increased to 40 h, the size of particles became nearly stable. Welding, fracturing and re-welding of the elemental powders occurred repeatedly during mechanical alloying process (Fig. 2(c) and (d)). From the SEM images with higher magnification (Fig. 2(e) and (f)), it can be
seen that the particles consisted of many nano-sized grains, and the size of the gains was decreased with the increase of the milling time. However, as shown in Table 1, considerable amount of impurities can be introduced if too long milling time is applied, in particular of Fe, a common occurrence during ball milling [13]. It has been found that high level of impurities had significant negative influence on the properties of W-Cu alloys, especially for thermal conductivity [8,9,17]. Therefore, the MA W-Cu powders milled for 40 h, with relatively low impurities level, were chosen to fabricate W-Cu alloys by HPS.
Table 1 Chemical analysis of MA W-Cu powders milled for different periods.
3.2. Densification results of HPS compacts
Milling time/hour
W
Cu
Fe
Cr
Ni
(wt%) Mn
5 15 40 60 80 100
82.13 78.42 80.03 79.93 79.98 80.03
17.81 21.48 19.79 19.71 19.51 19.14
0.033 0.063 0.118 0.263 0.336 0.574
0.019 0.025 0.048 0.084 0.153 0.231
0.004 0.006 0.007 0.009 0.012 0.018
0.001 0.002 0.004 0.006 0.007 0.009
The effects of temperature and pressure on the density and hardness of the sintered compacts were studied. Fig. 3(a) displays the densities and hardness of the compacts prepared by HPS of NS MA W-Cu powders at different temperatures under a pressure of 50 MPa. The densities are normalized to the theoretical density of W-20 wt% Cu of 15.64 g/cm3. As can be seen in Fig. 3(a), the densification of the compact was strongly influenced by the sintering temperature. From a density of 85.11%TD obtained at a sintering temperature of 900 °C, the density increased
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100
1100 oC
(b)
95 300 90 250
Hardness (HV)
Relative density (%TD)
350
relative density hardness
(a)
(c)
85 200 80
pores 150
75 950 1000 1050 o Sintering temperature ( C)
1100
350
100
Relative density (%TD)
(d)
relative density hardness
95 300 90 250
Hardness (HV)
900
85 200 80
150
75 0
20
30 Pressure (MPa)
40
50
Fig. 3. The bar graph illustrations for densities and hardness of the W-Cu sintered compacts prepared by HPS of NS MA W-Cu powders at different temperatures under a pressure of 50 MPa (a); the macroscopic photo of the die (b) and the SEM image of the W-Cu sintered compact (c) when HPS was carried out at 1100 °C under a pressure of 50 MPa; densities and hardness of the W-Cu sintered compacts prepared by HPS of NS MA W-Cu powders under different pressures at 1050 °C (d).
with sintering temperature, and reached the highest value (99.1%TD) at a sintering temperature of 1050 °C. An attempt to increase the sintering temperature into the liquid phase led to the loss of Cu melt out of the die (Fig. 3(b)), and to a high porosity within the samples (Fig. 3(c)). Fig. 3(d) displayed the densities and hardness of the compacts prepared by HPS of NS MA W-Cu powders under different pressures at
(a)
1050 °C. As shown, the pressure played an important role in the densification of compacts during HPS. The density of compact was only 86.2%TD without pressure (P = 0 MPa), and rose to 92.4%TD when a pressure of 20 MPa was applied. It reached 99.2%TD under a pressure of 40 MPa, and then remained nearly stable when the pressure was further increased.
(b)
copper pools
pores 10µm
10µm
Fig. 4. SEM images of W-Cu sintered compacts prepared from NS MA W-Cu powders (a) and standard mixed W-Cu powders (b).
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95
3.3. Properties of W-Cu alloy
100 hours CPS
The properties of the W-Cu alloy sintered at 1050 °C under a pressure of 40 MPa were investigated. Figs. 4(a) and (b) show SEM images of W-Cu sintered compacts prepared from NS MA W-Cu powders and that of the reference sample obtained from mixed W-Cu powders respectively. In Fig. 4, the light grey area is the W phase, while Cu appears as dark grey areas, and the pores appear as black areas. Compared to the compact prepared from mixed W-Cu powders, the compact prepared form NS MA W-Cu powders exhibited fewer pores and a more homogeneous microstructure consisting of ultra-fine W grains and a continuous Cu network, and some Cu rich pools outside the W particles (marked by the red arrows). Consistent with the morphologies shown in Fig. 4, the density of the compact prepared from NS MA W-Cu powder (99.2%TD) was much higher than that achieved using the mixed W-Cu powder (90%TD). After removal of Cu, the specimen prepared from mixed W-Cu powders collapsed, while the specimen prepared from NS MA W-Cu powders retained its shape. A SEM image of this sample is shown in Fig. 5. Although the sintering temperature was relatively low regarding W, sintering necks between W grains can be observed. This is likely to be the reason for the mechanical stability of the sample. Because of the low sintering temperature, the W grains coarsened slightly and had submicron size (about 500 nm). Consistent with the low porosity, the average hardness of the W-Cu sintered compacts prepared from NS MA powders reached close to 310 HV, much higher than that prepared from mixed W and Cu powders (average 260 HV).
80 hours 60 hours 40 hours 15 hours 5 hours W Cu
45
4.1. The effect of mechanical alloying process on the sintering properties of powders Fig. 6 shows the XRD patterns of MA W-Cu powders milled for different periods. With the increase of milling time, intensities of the diffraction peaks of Cu decreased rapidly and almost disappeared after 40 h milling. After milling for 100 h, only W diffraction peaks can be observed superposed on a broad background. A progressive broadening of the W peaks was observed with the increase of the milling time, which could be induced by grain refinement and lattice strain accumulation [16]. Generally, there are five factors contributing to the broadening of peaks appearing in XRD patterns: (1) crystallite size; (2) micro-strain; (3) diffraction instrument; (4) the overlapping of two sets of diffraction profiles from Kα1 and Kα2 diffraction of Cu; (5) scattering of background [18]. A standard Si sample was used to deal with the intrinsic instrument broadening effect. The effect of Cu Kα2 radiation on the peak was eliminated by Fourier series separation, and the scattering of background could be evaluated by the quantic polynomial expression.
(a)
W
60
75 o 2 ( )
W
W
90
105
W
120
Fig. 6. XRD patterns of MA W-Cu powders milled for different periods.
Then the crystallites size and lattice micro-strain can be related to the width of peaks by Williamson–Hall method [19]: β cosθ ¼
4. Discussions
Cu
W
Kλ þ 2ε sinθ d
ð1Þ
Where β is the full width at half-maximum (FWHM) of the MA powders, θ is the position of peak maximum, K is the Scherrer constant (about 0.9), λ is the radiation wavelength (about 1.5405 Å), d is the crystallite size and ε is the lattice micro-strain introduced by milling. Fig. 7 showed the variations of calculated crystallite size and microstrain of the W matrix with different milling times. The average crystallite size dropped rapidly in the first 5 h, and then dropped gradually to about 10.9 nm when the milling time increased to 40 h. After 40 h, the average crystallite sizes kept nearly constant, and the finest average crystallite size of 7.1 nm was attained at 100 h. In contrast to the variation of crystallite size, the micro-strain rose rapidly at the initial stage, and then increased gently after 40 h milling. As shown above, the MA W-Cu powders milled for 40 h exhibited a nano-crystalline microstructure and a large micro-strain, which obviously enhances the sintering properties of the powders [20]. We hold this is the major reason that the density of the sintered compact prepared from NS MA W-Cu powders is much higher than that prepared by simply mixed W-Cu powders.
(b)
10µm Fig. 5. SEM images of W skeleton after removal of Cu of the simple prepared from NS MA W-Cu powders.
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0.7
100
300
0.6
80
0.5
250
Crystallite size Micro-strain
0.4
200 0.3 150 0.2
Heat flow (W/g)
350
Micro-strain (%)
Crystallite size (nm)
96
60 40 20 0
100
1083 oC
0.1
50
-20
0.0
Exothermic -40
-0.1 0
20
40 60 Milling time (hour)
80
100
200
400
600
800
1000
Temperature (oC)
Fig. 7. Effects of milling time on crystallite size and micro-strain of the W matrix. Fig. 8. DTA curve of NS MA W-Cu powders at a heating rate of 10 °C/minute.
4.2. Densification mechanism of NS MA W-Cu powders during high pressure sintering The densification mechanism of NS W-Cu powders prepared by mechanical alloying process was different from that of the standard mixed W-Cu powders because of their different characteristic morphologies. As has been discussed in previous studies [20–28], the densification mechanism of NS MA W-Cu powders during sintering can be described as follows: the Cu-pools and Cu-layers are formed by the growth of Wgrains inside MA powders at the temperature near the melting point of Cu, referred to as “nanosintering phenomenon”, and the Cu phase diffuses out to the surface of individual MA powder particles and surrounds the exterior of MA powders, leading to a bonding between the particles. Further increase in density requires the rearrangement of particles. To induce this, increasing the temperature beyond the melting point of Cu is the most effective method, as the MA powder particles can be rearranged by the capillary force of liquid of the preformed Cupools and Cu-layers. It was reported that about 95%TD can be achieved by the particle rearrangement during liquid phase-sintering [21]. And Heady and Cahn [29] proposed that the maximum interparticle force due to capillary action was roughly equivalent to an external pressure: pffiffiffi P ¼ 2 2π ðγlv =RÞ cosθ
ð2Þ
Where γlv denotes the surface free energy of the liquid-vapor, R is radius of the particles, θ is the contact angle of liquid-solid, and P represents the equivalent external pressure. The higher the temperature, the smaller the θ. Therefore, full densification of standard mixed W-Cu powders was attained at temperatures above 1350 °C because a complete wetting of W-Cu system was then obtained, whereas full densification of NS MA W-Cu powders was already attained at about 1200 °C [21,22]. Fig. 8 showed the differential thermal analysis (DTA) image of NS MA W-Cu powders milled for 40 h. It can be seen that the melting point of Cu in the powders was about 1083 °C. The HPS process of NS MA powders was carried out at 1050 °C, below the melting point of Cu. Hence, the rearrangement of MA powder particles by capillary force could not be fully obtained. However, when an external pressure of 20 MPa was applied, the density of compact rose from 86.2%TD to 92.4%TD, and nearly full density (99.2%TD) could be obtained when the pressure was increased to 40 MPa, indicating that the rearrangement of MA powder particles can be driven by the external high pressure. The role of high pressure played in the densification of MA powders during sintering is equally important as the capillary force in liquid phase sintering. As a result, nearly full densification of MA powders can be achieved already at relatively low sintering temperature
under high pressure on the compact. Meanwhile, the temperature for HPS should be not too low. When the sintering temperature of HPS was much lower than the melting point of Cu, i.e., 950 °C, the density was still lower than 90%TD even at a pressure of 50 MPa (as shown in Fig. 4(a)). This we attribute to the fact that the rearrangement of MA powder particles is hindered in the absence of Cu-pools and Cu-layers which are usually formed at 1000 °C [21]. When the sintering temperature increased to 1050 °C, not only the Cu pools and layers were formed, but also the Cu phase became viscous and attained a certain extent of mobility at the temperature near the melting point of Cu [24]. Therefore, the rearrangement of MA powder particles was facilitated, and the gaps between MA powders could be filled with Cu phase when a pressure of 40 MPa was applied, resulting in nearly full density of the sintered compact. Because of the relatively low sintering temperature, the average size of W grains was about 500 nm, which was smaller than that of W grains sintered at the temperature of liquid phase (nearly or more than 1 μm) [21–23]. According to the study by Ryu and Kim, the bonding of MA powders sintered below the melting point of Cu was formed by the diffusion of Cu phase [20]. However, the bonding of MA powders during HPS was accomplished also by the sintering of W grains, as described in previous section. There are three main concerns on sintering of W– Cu nanocomposites: Complete densification, maintenance of nanocrystalline structure and shape retention [30,31]. Compared with liquid phase sintering, HPS is more conducive to achieve nanocrystalline structure and shape retention with nearly full densification. 5. Conclusions W-20 wt% Cu alloy with a high density (99.2%TD) and ultra-fine W grains (500 nm) can be successfully prepared by the HPS at 1050 °C under a pressure of 40 MPa. The changes of morphology and structural evolution of W-Cu powders during mechanical alloying process were observed and studied, and the effects of temperatures and pressure on the density of sintered compacts made from NS MA W-Cu powders were investigated. Based on these studies, the following conclusions can be made: 1. Nanosized MA W-Cu powders can be obtained after 40 h milling at a milling speed of 400 rpm. The results of XRD analysis indicate that these MA powders are nanocrystalline with a grain size around 10 nm. Extending the milling time did not reduce the W grain size significantly, but increased the impurity content, in particular of Fe, in the MA W-Cu powders. 2. The optimum temperature and pressure for HPS of NS MA W-Cu powders were 1050 °C and 40 MPa, respectively. At this temperature,
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ultra-fine W grains can be obtained in the final samples. The densification of MA powders could be strongly promoted due to sintering under pressure. The density increased with pressure, going into a limit at about 99%TD. 3. The densification of NS MA W-Cu powders during HPS at a temperature near the melting point of Cu is attributed to a nanosintering process accompanied by the rearrangement of MA powder particles under external high pressure. Acknowledgements The authors gratefully acknowledge financial support from the Natural Science Foundation of China (51271203), Self-dependent Innovation Project for Post-graduate Students of Central South University (2012zzts006), and grants from the project of innovation-driven plan in Central South University.
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