- Email: [email protected]
Confined fracture behavior of bulk metallic glass-coated tungsten composite wires produced by continuously coating process
Journal of Alloys and Compounds 553 (2013) 14–18
Contents lists available at SciVerse ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Confined fracture behavior of bulk metallic glass-coated tungsten composite wires produced by continuously coating process Baoyu Zhang, Xiaohua Chen ⇑, Xidong Hui ⇑ State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 10083, China
a r t i c l e
i n f o
Article history: Received 27 September 2012 Accepted 1 November 2012 Available online 10 November 2012 Keywords: Metallic glasses (MG) Coating material Microstructure Mechanical properties
a b s t r a c t The effects of thermal residual stresses on the tensile fracture behavior of the bulk metallic glass (BMG)coated composite wires have been investigated by fabricating a series of BMG composite wires at varies drawing velocity. It is found that the coating thickness increases with the increase of drawing velocity and the axial and radial thermal stresses of the composite wires increase with the increase of the coating thickness. The values of axial thermal stresses are comparable with the tensile strength difference between the composite wires and the tungsten wire. Due to the effects of radial thermal stresses, the fracture mode change from the unconfined cleavage fracture of pure tungsten wire to confined step-like fracture mode of composite wires. Ó 2012 Elsevier B.V. All rights reserved.
1. Introduction BMGs have become a hot topic of materials research because of their excellent mechanical, physical and chemical properties [1–3]. On one hand, amorphous microstructure leads to high strength [4], low modulus [5], high fracture toughness [6,7] and high fatigue limit [8–10]. On the other hand, no dislocation, grain boundary and other second phases make BMGs exhibit more improved corrosion resistance than their crystalline counterparts. A lot of work has been done on studying the corrosion behaviors of BMGs system in different aqueous solutions [11–16]. However, most of work only focused on monolithic BMGs. Meanwhile, BMGs as one of the best coating material are needed to be coated on the metallic wire for protection. BMG-coated wire composites have been seldom reported [17,18]. Applying BMGs as coating material also can overcome the disadvantage of size limitation of BMGs. So, it is significant to study the synthesizing process and mechanical properties of BMG-coated wire composite for aggressive environment application. To obtain BMG coating, we developed a method [19,20], as shown in Fig. 1, which presents the schematic of continuous preparation process of BMG-coated composite wires. The continuous preparation unit consists of a vacuum unit and a crucible for a molten metal surrounded with the heater. The wire bundles that are fed continuously from a supply wheel must first pass through preheating unit, and then immersed into molten metal for infiltration. After infiltration the wire bundles pass through the argon cooling unit that offer an enough cooling rate to successfully produce ⇑ Corresponding authors. Tel./fax: +86 010 62332350. E-mail addresses: [email protected] (X. Chen), [email protected] (X. Hui). 0925-8388/$ - see front matter Ó 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2012.11.002
BMG coating and finally exit to a take-up wheel. It was found that Vit1 BMG-coated tungsten composite wires produced by this continuous method possessed similar ultimate strength and ductility as the bare tungsten wire under tensile loading [20]. The processing parameters such as the speed of the wire movement, infiltration time and melt temperature can be controlled by this unit through setting a series of controlling programs, respectively. In this paper, effects of thermal residual stresses on the fracture behaviors of composite wires were investigated. 2. Experimental procedures Pre-alloy ingots of Zr41Ti14Cu12.5Ni10Be22.5 (Vit1, at.%) were prepared by arc melting elements with high purity under a Ti-gettered argon atmosphere. The ingots were remelted four times for homogenizing the alloy elements. Tungsten wires with diameter of 500 lm were polished by an electrochemical method [21] and cleaned in an ultrasonic bat of acetone and ethanol. Alloy ingots were placed in a crucible and heated to a coating temperature measured using a far infrared thermometer. By using continuous production unit sketched in Fig. 1, composite wires obtained with drawing velocities of 10, 30, 40, and 50 mm/s, respectively, at 1023 K, were produced successfully (see the details in reference 19). X-ray diffraction (XRD) with Cu Ka radiation was done using the Rigaku D/max-rB X-ray diffractometer. Differential scanning calorimeter (DSC) was performed using Netzsch STA449 calorimeter with a heating rate of 20 K/min. Tensile properties of BMG-coated composite wires at room temperature were tested on Reger RGM 3010 with the strain rate of 1 104 s1. The microstructural observation and fracture surface of BMG-coated composite wires were carried out by scanning electron microscopy (SEM) using a Carl Zeiss Auriga Crossbeam Workstation.
3. Results and discussion Fig. 2a shows the XRD patterns of BMG-coated composite wires and uncoated tungsten wire and monolithic BMG Vit1, respectively. The patterns of composite wires show some sharp crystalline
15
B. Zhang et al. / Journal of Alloys and Compounds 553 (2013) 14–18
Fig. 1. Schematic diagram of continuously coating unit.
Intensity (a.u.)
(a)
(110)
W wire
(210)
(211)
(220)
50 mm/s 40 mm/s 30 mm/s 10 mm/s Vit 1 20
40
60
80
100
2 theta (degree)
Endothermic flow (mW/mg)
(b)
500
Tg
50 mm/s
Tg
40 mm/s
Tg
30 mm/s
Tg
10 mm/s BMG Vit1 Tg
600
Tx Tx Tx Tx Tx
700
800
Temperature (K) Fig. 2. XRD patterns and DSC curves of BMG-coated composite wires fabricated with various drawing velocities at 1023 K. Data of pure tungsten and monolithic BMG are presented for comparison.
diffraction peaks of tungsten crystalline superimposed on the broad scattering hump which is typical characteristic of amorphous structure. So the coating is mainly comprised of amorphous phase. Meanwhile, some other crystalline diffraction peaks are detected at 10 mm/s and 1023 K. They are indexed as phase of Ni2Zr3, which indicate the decrease of glass forming ability (GFA) as a result of the diffusion of tungsten elements [22]. It can be noted that the intensity of diffraction peaks corresponding to the tungsten phase declines with the increase of the drawing velocity, implying that the BMG coating layer becomes thicker at a higher drawing velocity.
The DSC curves of BMG-coated composite wires and monolithic BMG vit1 are shown in Fig. 2b. The glass transition temperature (Tg), the onset crystallization temperature (Tx1) are marked by arrows. It can be noted that all curves show the characteristics of a glass transition endothermic peak followed by multiple crystallization exothermic peaks. The data of Tg (629–637 K) and Tx1 (692– 702 K) of BMG coating also agree with the data of monolithic BMG samples and the reported data by other researchers [23]. So the amorphous structure of coating can be confirmed by the results of XRD and DSC. Fig. 3 shows energy selective backscattering (ESB) images of the cross-section microstructure of the composite wires, demonstrating the outer coating layer (at the top) and the inner tungsten matrix (at the bottom) at a drawing velocity of (a) 10 mm/s, (b)30 mm/s, (c) 40 mm/s, and (d) 50 mm/s at 1023 K, respectively. As shown, uniform contrast was observed at the interface between the BMG coating and the tungsten matrix, which are consistent with the XRD pattern in Fig. 1b. coating thickness at various drawing velocities are presented in Table 1. It is clear that the coating thickness increases with increase of the drawing velocity. Fig. 3e shows the dependence of drawing velocity and coating temperature on BMG coating thickness. The amorphous coating thickness increases with the increasing drawing velocity, which roughly obeys a parabolic relationship. By fitting the curve, the relationship between coating thickness (d) and drawing velocity (m) can be approximately expressed as d = 2.35 + 0.10 m 6.52 m2. Thus, by controlling the drawing velocity, different thickness of coating can be achieved. The increase of drawing velocity causes the increase of the cooling rate, which will suppress the nucleation of the crystal phase and form amorphous structure. The results agree with the results of XRD patterns and DSC curves in Fig. 2a and b. The tensile engineering stress–strain curves of BMG-coated composite wires are shown in Fig. 4 along with pure tungsten wire and monolithic Vit1 BMG for comparison. All composite wires exhibit the characteristics of obvious macroscopic plasticity and work-hardening. It is observed that they all show similar ultimate tensile strength to the pure tungsten wire. But, with 30 mm/s, the composite wire shows higher tensile strength and bigger plasticity than the pure tungsten wire. The composite wire obtained at 30 mm/s shows the best comprehensive mechanical properties with ultimate tensile strength of 2335 MPa and tensile plasticity strain of 1.52%, respectively. Compared with the tungsten wire, it is noticeable that the tensile strength and plasticity of the composite wire with 30 mm/sis increased by 25 MPa and 0.25%, respectively. Due to the mismatch of the coefficients of the thermal expansion (CTE) between the tungsten wire and the Vit1 coating, the thermal residual stresses exist in the composite wires when cooling from high temperature of 1023 K to room temperature. The tensile fracture behavior of the composite wire is affected by the stress concentration at the interface because of thermal residual stresses. So, it is necessary to investigate the effect of thermal residual stress on the facture behavior of the composite wire. The BMG-coated tungsten composite wire is modeled as a perfectly bonded coaxial composite cylinder, as shown in Fig. 5a and b. The classical Lame solution is used to calculate the principal stresses in thick cylinders. The axial and radial thermal stresses at the interface with different coating thickness are calculated using the method illustrated by Choy [24]:
rr ¼
rh
a22 d1 a12 d2 r2 1 c2 D rf
a22 d1 a12 d2 r2 ¼ 1 c2 D rf
! ð1Þ ! ð2Þ
16
B. Zhang et al. / Journal of Alloys and Compounds 553 (2013) 14–18
(a)
(b)
µ 4 µm
µ 4 µm (c)
(d)
µm 4µ
µ 4 µm
Fig. 3. Cross-sectional energy selective backscattering (ESB) images of BMG-coated composite wires fabricated with drawing velocities of 10 mm/s (a), 30 mm/s (b), and 50 mm/s (c), respectively.
Table 1 BMG coating thickness at drawing velocities of 10 mm/s, 30 mm/s, 40 mm/s and 50 mm/s for BMG-coated composite wires. Drawing velocity (mm/s) Mean BMG coating thickness (lm)
10 3.32
30 4.78
40 5.50
(a)
(b)
50 5.83
2500
20
axial residual stress
Thermal resudual stress (MPa)
Engineering Stress (MPa)
(c) 2000
1500
1000
10 mm/s 30 mm/s 40 mm/s 50 mm/s tungsten wire
500
0
1
2
3
4
5
experimental value of
z r w
-
c
0
-20
-40
-60
Engineering Strain (%)
radial residual stress
0
2
4
6
8
10
Coating thickness ( m) Fig. 4. Uniaxial tensile stress–strain curves of BMG-coated composite wires and pure tungsten wire, respectively.
rz
a11 d2 a21 d1 r2 ¼ 1 c2 D rf
! ð3Þ
where rc and rf are the external coating radius and the tungsten wire radius, respectively; the constants of a11, a22, a12, a21, d1, d2 and D are given explicitly in term of the thermoelastic properties of the tungsten wire and the coating, which are as followed [24]:
a11 ¼ 2a22
a12
2cf 2c ¼ cþ Ec Ef
1 1 ¼ Ec Ef
r2 1 c2 rf
r2 1 c2 rf
! ð4Þ
! ð5Þ
Fig. 5. Thermal residual stress calculation model for BMG-coated composite wires with external coating radius rc and the tungsten wire radius rf (a, b), and coating thickness dependence of the axial thermal residual stress rz, radial thermal residual stress rr and tensile strength differences between the tungsten wire rw and BMGcoated composite wires rc (c), respectively.
a21
! 1 cf 1 cc r 2c 1 þ cc r 2c 1 2 ¼ þ 2 Ec Ec Ef rf rf
d1 ¼ d2 ¼ ðac af ÞðT 1 T 2 Þ
ð6Þ ð7Þ
where cc and cf are the Poisson’s ratio of the coating and the tungsten wire, respectively; Ec and Ef are the Young’s modulus of the coating and the tungsten wire, respectively; ac and af are the coefficient of thermal expansion (CTE) of the coating and the tungsten wire, respectively; T1 and T2 are the manufacturing temperature and the room temperature, respectively. The data used in
B. Zhang et al. / Journal of Alloys and Compounds 553 (2013) 14–18
thermal residual stress calculation are presented as the following: for BMG, Ec = 96 GPa [25], cc = 0.36 [25], and ac = 13 106 K1 (ac = 9.0–15 106 K1 [26]); for tungsten wire, Ef = 410 GPa [27], cf = 0.28 [25], and af = 4.6 106 k1 (af = 4.5–4.7 106 K1 [28]). As shown in Fig. 5c, it can be seen that the axial and radial thermal stress loading on the tungsten wire are all negative value, which indicate that both are compressive stress. The axial thermal stresses increase quickly while the radial thermal stresses increase slowly. The existence of axial stress and radial stress will restrict the propagation of the radial cracks, leading the radial cracks tilting to form axial cracks [29]. It can be noticed that the value of the axial thermal stress is comparable with that of stress difference between the composite wire and the tungsten wire at 30 mm/s, 40 mm/s and 50 mm/s, respectively, while it shows a big deviation from the calculated value at 10 mm/s. It can be explained from the Fig. 2a that the phase of Ni2Zr3 acts as weak sites to leading early failure when deformation. So, it is reasonable to think that the tensile strength of composite wire increase partly due to confined effect of relative big axial thermal stress. Before exploring the fracture mechanism of the composite wires, we will have a look at tensile fracture surface of pure tungsten wire. Fig. 6a shows the tensile fracture morphology of pure tungsten wire after tensile test. As shown in Fig. 6a there are two different features: at the central zone is smooth fracture plane, which is the typical fracture feature of the transgranular fracture and at the edge zone are the intergranular fracture zone with characteristics of radial cracks and some deep axial cracks initiating from the surface cracks as shown in Fig. 6b while necking. When
17
deformation, the cracks nucleate in the inner defect zone of the tungsten fiber and propagate quickly towards the edge zone by the way of cleavage fracture, leading to fracture nearly without forming shear offset zone. The axial cracks form as a result of necking, which is the cause of the tensile plasticity of tungsten fiber. Fig. 6b shows the side view of the tungsten wire after fracture. It can be seen that the fracture plane is nearly perpendicular with the axial direction of the tungsten wire, indicating the cracks’ propagation without confining. The fracture surface of BMG-coated composite wire and its coating synthesized at 30 mm/s are shown in Fig. 6c and d, respectively. The cracks firstly initiate from the central zone of the tungsten wire and spread towards the surrounding coating. Fig. 6d shows the side morphology of BMGcoated composite wire. Due to the confined effect of the radial thermal stress, the spread of axial cracks tilt to form the transverse macro-cracks, leading to step-like pattern fracture. Fig. 6e shows the river-like pattern and the tearing of the shear bands at the intersection site of shear bands, indicating the amorphous coating undergoing severe plastic deformation. The white arrows show the shear step, and black arrows designate the local melting site. It can be also shown from Fig. 6e that the dense microcracks locate in the side of tungsten wire, which indicates the good interface bonding between the coating and the tungsten fiber. As shown in Fig. 6f, dense patterns of primary and secondary shear bands as well as the microcracks along the shear bands are visible, which is consistent with the improving tensile strength and plasticity of the composite. Compared Fig. 6b with Fig. 6d, due to the existence of thermal residual stress, the fracture mode of tungsten wire
(a)
(b)
(c)
(d)
(e)
(f)
Fig. 6. SEM images of fracture surfaces and their side view images: (a, b) the pure tungsten wire; (c, d) BMG-coated composite wires at 30 mm/s; (e, f) the magnified images corresponding to (c) and (d), respectively.
18
B. Zhang et al. / Journal of Alloys and Compounds 553 (2013) 14–18
changes from cleavage fracture mode of pure tungsten to step-like fracture mode of composite wire, which is the cause of improving tensile strength and plasticity for composite wires. 4. Conclusions At drawing velocity of 30 mm/s and processing temperate of 1023 K, the composite wires show the best comprehensive mechanical properties with tensile strength of 2235 MPa and plasticity of 1.52%. The fracture behavior of BMG-coated composite wires is determined by the thermal residual stresses. Due to the confined effects of axial stresses and radial stresses, the propagation of the radial cracks is restricted to tilt, forming the axial cracks. So the fracture mode changes from unconfined cleavage fracture of pure tungsten wire to step-like fracture of BMG-coated composite wires. Acknowledgments The authors are grateful for the financial support of the Specialized Research Fund for the Doctoral Program of Higher Education of China (No. 20100006120020) and National Nature Science Foundation of China (51071018 and 51271018). The authors also thank the support of the Carl Zeiss Auriga Crossbeam Workstation in the State Key Laboratory of Advanced Metals and Materials. References [1] [2] [3] [4] [5]
A. Inoue, Acta Mater. 48 (2000) 279–306. W.H. Wang, C. Dong, C.H. Shek, Mater. Sci. Eng. R 44 (2004) 45–89. A.L. Greer, E. Ma, MRS Bull. 32 (2007) 611–619. J.F. Wang, R. Li, N.B. Hua, T. Zhang, J. Mater. Res. 26 (2011) 2072–2079. D.G. Pan, W.Y. Liu, H.F. Zhang, A.M. Wang, Z.Q. Hu, J. Alloys Comp. 438 (2007) 142–144.
[6] M.D. Demetriou, M.E. Launey, G. Garrett, J.P. Schramm, D.C. Hofmann, W.L. Johnson, R.O. Ritchie, Nat. Mater. 10 (2011) 123–128. [7] Q. He, Y.Q. Cheng, E. Ma, J. Xu, Acta Mater. 59 (2011) 202–215. [8] M.E. Launey, D.C. Hofmann, W.L. Johnson, R.O. Ritchie, Proc. Natl. Acad. Sci. USA 106 (2009) 4986–4991. [9] D.C. Qiao, G.J. Fan, P.K. Liaw, H. Choo, Int. J. Fatigue 29 (2007) 2149–2154. [10] Y. Nakai, S. Hosomi, Mater. Trans. JIM 48 (2007) 1770–1773. [11] L. Huang, D.C. Qiao, B.A. Green, P.K. Liaw, J.F. Wang, S.J. Pang, T. Zhang, Intermetallics 17 (2009) 195–199. [12] U.K. Mudali, S. Baunack, J. Eckert, L. Schultz, A. Gebert, J. Alloys Comp. 377 (2004) 290–297. [13] K. Asami, C.L. Qin, T. Zhang, A. Inoue, Mater. Sci. Eng. A 375–377 (2004) 235– 239. [14] C.L. Qin, K. Asami, T. Zhang, W. Zhang, A. Inoue, Mater. Trans. JIM 44 (2003) 749–753. [15] V.R. Raju, U. Kühn, U. Wolff, F. Schneider, J. Eckert, R. Reiche, A. Gebert, Mater. Lett. 57 (2002) 173–177. [16] W.H. Peter, R.A. Buchanan, C.T. Liu, P.K. Liaw, M.L. Morrison, J.A. Horton, C.A. Carmichael Jr., J.L. Wright, Intermetallics 10 (2002) 1157–1162. [17] X.H. Chen, B.Y. Zhang, G.L. Chen, X.D. Hui, X.R. Xing, Intermetallics 19 (2011) 1913–1918. [18] P. Yu, K.C. Chan, L. Xia, H.B. Yu, H.Y. Bai, Mater. Trans. JIM 50 (2009) 2451– 2454. [19] Y. Zhang, X.H. Chen, G.L. Chen, X.C. Zhang, Z.D. Wang, China patent ZL200710120355.4, 2009. [20] X.H. Chen, B.Y. Zhang, G.L. Chen, Y. Zhang, X.D. Hui, Z.P. Lu, X.J. Liu, Y. Xu, X.R. Xing, Intermetallics 18 (2010) 2034–2038. [21] X.H. Chen, Y. Zhang, X.C. Zhang, X.D. Hui, B.C. Wei, G.L. Chen, Electrochem. Solid-State Lett. 10 (2007) 21–23. [22] M.L. Wang, G.L. Chen, X. Hui, Y. Zhang, Z.Y. Bai, Intermetallics 15 (2007) 1309– 1315. [23] A. Peker, W.L. Johnson, Appl. Phys. Lett. 63 (1993) 2342–2344. [24] K.L. Choy, J.F. Durodola, B. Derby, C. Ruiz, Composites 26 (1995) 531–539. [25] R.D. Conner, R.B. Dandliker, W.L. Johnson, Acta Mater. 46 (1998) 6089–6102. [26] Y. He, R.B. Schwarz, D.G. Mandrus, J. Mater. Res. 11 (1996) 1836–1841. [27] H.E. Boyer, T.L. Gall, Metals Handbook, Desk Edition., American Society for Metals, Metals Park, Ohio, 1985. [28] Y.S. Touloukian, R.K. Kirby, R.E. Taylor, P.D. Desai, Thermal Expansion: Metallic Elements and Alloys, IFI/Plenum, New York, 1975. [29] J. Zarzycki, Glasses and Vitreous State, Cambridge University Press, Cambridge, 1991.
Contents lists available at SciVerse ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Confined fracture behavior of bulk metallic glass-coated tungsten composite wires produced by continuously coating process Baoyu Zhang, Xiaohua Chen ⇑, Xidong Hui ⇑ State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 10083, China
a r t i c l e
i n f o
Article history: Received 27 September 2012 Accepted 1 November 2012 Available online 10 November 2012 Keywords: Metallic glasses (MG) Coating material Microstructure Mechanical properties
a b s t r a c t The effects of thermal residual stresses on the tensile fracture behavior of the bulk metallic glass (BMG)coated composite wires have been investigated by fabricating a series of BMG composite wires at varies drawing velocity. It is found that the coating thickness increases with the increase of drawing velocity and the axial and radial thermal stresses of the composite wires increase with the increase of the coating thickness. The values of axial thermal stresses are comparable with the tensile strength difference between the composite wires and the tungsten wire. Due to the effects of radial thermal stresses, the fracture mode change from the unconfined cleavage fracture of pure tungsten wire to confined step-like fracture mode of composite wires. Ó 2012 Elsevier B.V. All rights reserved.
1. Introduction BMGs have become a hot topic of materials research because of their excellent mechanical, physical and chemical properties [1–3]. On one hand, amorphous microstructure leads to high strength [4], low modulus [5], high fracture toughness [6,7] and high fatigue limit [8–10]. On the other hand, no dislocation, grain boundary and other second phases make BMGs exhibit more improved corrosion resistance than their crystalline counterparts. A lot of work has been done on studying the corrosion behaviors of BMGs system in different aqueous solutions [11–16]. However, most of work only focused on monolithic BMGs. Meanwhile, BMGs as one of the best coating material are needed to be coated on the metallic wire for protection. BMG-coated wire composites have been seldom reported [17,18]. Applying BMGs as coating material also can overcome the disadvantage of size limitation of BMGs. So, it is significant to study the synthesizing process and mechanical properties of BMG-coated wire composite for aggressive environment application. To obtain BMG coating, we developed a method [19,20], as shown in Fig. 1, which presents the schematic of continuous preparation process of BMG-coated composite wires. The continuous preparation unit consists of a vacuum unit and a crucible for a molten metal surrounded with the heater. The wire bundles that are fed continuously from a supply wheel must first pass through preheating unit, and then immersed into molten metal for infiltration. After infiltration the wire bundles pass through the argon cooling unit that offer an enough cooling rate to successfully produce ⇑ Corresponding authors. Tel./fax: +86 010 62332350. E-mail addresses: [email protected] (X. Chen), [email protected] (X. Hui). 0925-8388/$ - see front matter Ó 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2012.11.002
BMG coating and finally exit to a take-up wheel. It was found that Vit1 BMG-coated tungsten composite wires produced by this continuous method possessed similar ultimate strength and ductility as the bare tungsten wire under tensile loading [20]. The processing parameters such as the speed of the wire movement, infiltration time and melt temperature can be controlled by this unit through setting a series of controlling programs, respectively. In this paper, effects of thermal residual stresses on the fracture behaviors of composite wires were investigated. 2. Experimental procedures Pre-alloy ingots of Zr41Ti14Cu12.5Ni10Be22.5 (Vit1, at.%) were prepared by arc melting elements with high purity under a Ti-gettered argon atmosphere. The ingots were remelted four times for homogenizing the alloy elements. Tungsten wires with diameter of 500 lm were polished by an electrochemical method [21] and cleaned in an ultrasonic bat of acetone and ethanol. Alloy ingots were placed in a crucible and heated to a coating temperature measured using a far infrared thermometer. By using continuous production unit sketched in Fig. 1, composite wires obtained with drawing velocities of 10, 30, 40, and 50 mm/s, respectively, at 1023 K, were produced successfully (see the details in reference 19). X-ray diffraction (XRD) with Cu Ka radiation was done using the Rigaku D/max-rB X-ray diffractometer. Differential scanning calorimeter (DSC) was performed using Netzsch STA449 calorimeter with a heating rate of 20 K/min. Tensile properties of BMG-coated composite wires at room temperature were tested on Reger RGM 3010 with the strain rate of 1 104 s1. The microstructural observation and fracture surface of BMG-coated composite wires were carried out by scanning electron microscopy (SEM) using a Carl Zeiss Auriga Crossbeam Workstation.
3. Results and discussion Fig. 2a shows the XRD patterns of BMG-coated composite wires and uncoated tungsten wire and monolithic BMG Vit1, respectively. The patterns of composite wires show some sharp crystalline
15
B. Zhang et al. / Journal of Alloys and Compounds 553 (2013) 14–18
Fig. 1. Schematic diagram of continuously coating unit.
Intensity (a.u.)
(a)
(110)
W wire
(210)
(211)
(220)
50 mm/s 40 mm/s 30 mm/s 10 mm/s Vit 1 20
40
60
80
100
2 theta (degree)
Endothermic flow (mW/mg)
(b)
500
Tg
50 mm/s
Tg
40 mm/s
Tg
30 mm/s
Tg
10 mm/s BMG Vit1 Tg
600
Tx Tx Tx Tx Tx
700
800
Temperature (K) Fig. 2. XRD patterns and DSC curves of BMG-coated composite wires fabricated with various drawing velocities at 1023 K. Data of pure tungsten and monolithic BMG are presented for comparison.
diffraction peaks of tungsten crystalline superimposed on the broad scattering hump which is typical characteristic of amorphous structure. So the coating is mainly comprised of amorphous phase. Meanwhile, some other crystalline diffraction peaks are detected at 10 mm/s and 1023 K. They are indexed as phase of Ni2Zr3, which indicate the decrease of glass forming ability (GFA) as a result of the diffusion of tungsten elements [22]. It can be noted that the intensity of diffraction peaks corresponding to the tungsten phase declines with the increase of the drawing velocity, implying that the BMG coating layer becomes thicker at a higher drawing velocity.
The DSC curves of BMG-coated composite wires and monolithic BMG vit1 are shown in Fig. 2b. The glass transition temperature (Tg), the onset crystallization temperature (Tx1) are marked by arrows. It can be noted that all curves show the characteristics of a glass transition endothermic peak followed by multiple crystallization exothermic peaks. The data of Tg (629–637 K) and Tx1 (692– 702 K) of BMG coating also agree with the data of monolithic BMG samples and the reported data by other researchers [23]. So the amorphous structure of coating can be confirmed by the results of XRD and DSC. Fig. 3 shows energy selective backscattering (ESB) images of the cross-section microstructure of the composite wires, demonstrating the outer coating layer (at the top) and the inner tungsten matrix (at the bottom) at a drawing velocity of (a) 10 mm/s, (b)30 mm/s, (c) 40 mm/s, and (d) 50 mm/s at 1023 K, respectively. As shown, uniform contrast was observed at the interface between the BMG coating and the tungsten matrix, which are consistent with the XRD pattern in Fig. 1b. coating thickness at various drawing velocities are presented in Table 1. It is clear that the coating thickness increases with increase of the drawing velocity. Fig. 3e shows the dependence of drawing velocity and coating temperature on BMG coating thickness. The amorphous coating thickness increases with the increasing drawing velocity, which roughly obeys a parabolic relationship. By fitting the curve, the relationship between coating thickness (d) and drawing velocity (m) can be approximately expressed as d = 2.35 + 0.10 m 6.52 m2. Thus, by controlling the drawing velocity, different thickness of coating can be achieved. The increase of drawing velocity causes the increase of the cooling rate, which will suppress the nucleation of the crystal phase and form amorphous structure. The results agree with the results of XRD patterns and DSC curves in Fig. 2a and b. The tensile engineering stress–strain curves of BMG-coated composite wires are shown in Fig. 4 along with pure tungsten wire and monolithic Vit1 BMG for comparison. All composite wires exhibit the characteristics of obvious macroscopic plasticity and work-hardening. It is observed that they all show similar ultimate tensile strength to the pure tungsten wire. But, with 30 mm/s, the composite wire shows higher tensile strength and bigger plasticity than the pure tungsten wire. The composite wire obtained at 30 mm/s shows the best comprehensive mechanical properties with ultimate tensile strength of 2335 MPa and tensile plasticity strain of 1.52%, respectively. Compared with the tungsten wire, it is noticeable that the tensile strength and plasticity of the composite wire with 30 mm/sis increased by 25 MPa and 0.25%, respectively. Due to the mismatch of the coefficients of the thermal expansion (CTE) between the tungsten wire and the Vit1 coating, the thermal residual stresses exist in the composite wires when cooling from high temperature of 1023 K to room temperature. The tensile fracture behavior of the composite wire is affected by the stress concentration at the interface because of thermal residual stresses. So, it is necessary to investigate the effect of thermal residual stress on the facture behavior of the composite wire. The BMG-coated tungsten composite wire is modeled as a perfectly bonded coaxial composite cylinder, as shown in Fig. 5a and b. The classical Lame solution is used to calculate the principal stresses in thick cylinders. The axial and radial thermal stresses at the interface with different coating thickness are calculated using the method illustrated by Choy [24]:
rr ¼
rh
a22 d1 a12 d2 r2 1 c2 D rf
a22 d1 a12 d2 r2 ¼ 1 c2 D rf
! ð1Þ ! ð2Þ
16
B. Zhang et al. / Journal of Alloys and Compounds 553 (2013) 14–18
(a)
(b)
µ 4 µm
µ 4 µm (c)
(d)
µm 4µ
µ 4 µm
Fig. 3. Cross-sectional energy selective backscattering (ESB) images of BMG-coated composite wires fabricated with drawing velocities of 10 mm/s (a), 30 mm/s (b), and 50 mm/s (c), respectively.
Table 1 BMG coating thickness at drawing velocities of 10 mm/s, 30 mm/s, 40 mm/s and 50 mm/s for BMG-coated composite wires. Drawing velocity (mm/s) Mean BMG coating thickness (lm)
10 3.32
30 4.78
40 5.50
(a)
(b)
50 5.83
2500
20
axial residual stress
Thermal resudual stress (MPa)
Engineering Stress (MPa)
(c) 2000
1500
1000
10 mm/s 30 mm/s 40 mm/s 50 mm/s tungsten wire
500
0
1
2
3
4
5
experimental value of
z r w
-
c
0
-20
-40
-60
Engineering Strain (%)
radial residual stress
0
2
4
6
8
10
Coating thickness ( m) Fig. 4. Uniaxial tensile stress–strain curves of BMG-coated composite wires and pure tungsten wire, respectively.
rz
a11 d2 a21 d1 r2 ¼ 1 c2 D rf
! ð3Þ
where rc and rf are the external coating radius and the tungsten wire radius, respectively; the constants of a11, a22, a12, a21, d1, d2 and D are given explicitly in term of the thermoelastic properties of the tungsten wire and the coating, which are as followed [24]:
a11 ¼ 2a22
a12
2cf 2c ¼ cþ Ec Ef
1 1 ¼ Ec Ef
r2 1 c2 rf
r2 1 c2 rf
! ð4Þ
! ð5Þ
Fig. 5. Thermal residual stress calculation model for BMG-coated composite wires with external coating radius rc and the tungsten wire radius rf (a, b), and coating thickness dependence of the axial thermal residual stress rz, radial thermal residual stress rr and tensile strength differences between the tungsten wire rw and BMGcoated composite wires rc (c), respectively.
a21
! 1 cf 1 cc r 2c 1 þ cc r 2c 1 2 ¼ þ 2 Ec Ec Ef rf rf
d1 ¼ d2 ¼ ðac af ÞðT 1 T 2 Þ
ð6Þ ð7Þ
where cc and cf are the Poisson’s ratio of the coating and the tungsten wire, respectively; Ec and Ef are the Young’s modulus of the coating and the tungsten wire, respectively; ac and af are the coefficient of thermal expansion (CTE) of the coating and the tungsten wire, respectively; T1 and T2 are the manufacturing temperature and the room temperature, respectively. The data used in
B. Zhang et al. / Journal of Alloys and Compounds 553 (2013) 14–18
thermal residual stress calculation are presented as the following: for BMG, Ec = 96 GPa [25], cc = 0.36 [25], and ac = 13 106 K1 (ac = 9.0–15 106 K1 [26]); for tungsten wire, Ef = 410 GPa [27], cf = 0.28 [25], and af = 4.6 106 k1 (af = 4.5–4.7 106 K1 [28]). As shown in Fig. 5c, it can be seen that the axial and radial thermal stress loading on the tungsten wire are all negative value, which indicate that both are compressive stress. The axial thermal stresses increase quickly while the radial thermal stresses increase slowly. The existence of axial stress and radial stress will restrict the propagation of the radial cracks, leading the radial cracks tilting to form axial cracks [29]. It can be noticed that the value of the axial thermal stress is comparable with that of stress difference between the composite wire and the tungsten wire at 30 mm/s, 40 mm/s and 50 mm/s, respectively, while it shows a big deviation from the calculated value at 10 mm/s. It can be explained from the Fig. 2a that the phase of Ni2Zr3 acts as weak sites to leading early failure when deformation. So, it is reasonable to think that the tensile strength of composite wire increase partly due to confined effect of relative big axial thermal stress. Before exploring the fracture mechanism of the composite wires, we will have a look at tensile fracture surface of pure tungsten wire. Fig. 6a shows the tensile fracture morphology of pure tungsten wire after tensile test. As shown in Fig. 6a there are two different features: at the central zone is smooth fracture plane, which is the typical fracture feature of the transgranular fracture and at the edge zone are the intergranular fracture zone with characteristics of radial cracks and some deep axial cracks initiating from the surface cracks as shown in Fig. 6b while necking. When
17
deformation, the cracks nucleate in the inner defect zone of the tungsten fiber and propagate quickly towards the edge zone by the way of cleavage fracture, leading to fracture nearly without forming shear offset zone. The axial cracks form as a result of necking, which is the cause of the tensile plasticity of tungsten fiber. Fig. 6b shows the side view of the tungsten wire after fracture. It can be seen that the fracture plane is nearly perpendicular with the axial direction of the tungsten wire, indicating the cracks’ propagation without confining. The fracture surface of BMG-coated composite wire and its coating synthesized at 30 mm/s are shown in Fig. 6c and d, respectively. The cracks firstly initiate from the central zone of the tungsten wire and spread towards the surrounding coating. Fig. 6d shows the side morphology of BMGcoated composite wire. Due to the confined effect of the radial thermal stress, the spread of axial cracks tilt to form the transverse macro-cracks, leading to step-like pattern fracture. Fig. 6e shows the river-like pattern and the tearing of the shear bands at the intersection site of shear bands, indicating the amorphous coating undergoing severe plastic deformation. The white arrows show the shear step, and black arrows designate the local melting site. It can be also shown from Fig. 6e that the dense microcracks locate in the side of tungsten wire, which indicates the good interface bonding between the coating and the tungsten fiber. As shown in Fig. 6f, dense patterns of primary and secondary shear bands as well as the microcracks along the shear bands are visible, which is consistent with the improving tensile strength and plasticity of the composite. Compared Fig. 6b with Fig. 6d, due to the existence of thermal residual stress, the fracture mode of tungsten wire
(a)
(b)
(c)
(d)
(e)
(f)
Fig. 6. SEM images of fracture surfaces and their side view images: (a, b) the pure tungsten wire; (c, d) BMG-coated composite wires at 30 mm/s; (e, f) the magnified images corresponding to (c) and (d), respectively.
18
B. Zhang et al. / Journal of Alloys and Compounds 553 (2013) 14–18
changes from cleavage fracture mode of pure tungsten to step-like fracture mode of composite wire, which is the cause of improving tensile strength and plasticity for composite wires. 4. Conclusions At drawing velocity of 30 mm/s and processing temperate of 1023 K, the composite wires show the best comprehensive mechanical properties with tensile strength of 2235 MPa and plasticity of 1.52%. The fracture behavior of BMG-coated composite wires is determined by the thermal residual stresses. Due to the confined effects of axial stresses and radial stresses, the propagation of the radial cracks is restricted to tilt, forming the axial cracks. So the fracture mode changes from unconfined cleavage fracture of pure tungsten wire to step-like fracture of BMG-coated composite wires. Acknowledgments The authors are grateful for the financial support of the Specialized Research Fund for the Doctoral Program of Higher Education of China (No. 20100006120020) and National Nature Science Foundation of China (51071018 and 51271018). The authors also thank the support of the Carl Zeiss Auriga Crossbeam Workstation in the State Key Laboratory of Advanced Metals and Materials. References [1] [2] [3] [4] [5]
A. Inoue, Acta Mater. 48 (2000) 279–306. W.H. Wang, C. Dong, C.H. Shek, Mater. Sci. Eng. R 44 (2004) 45–89. A.L. Greer, E. Ma, MRS Bull. 32 (2007) 611–619. J.F. Wang, R. Li, N.B. Hua, T. Zhang, J. Mater. Res. 26 (2011) 2072–2079. D.G. Pan, W.Y. Liu, H.F. Zhang, A.M. Wang, Z.Q. Hu, J. Alloys Comp. 438 (2007) 142–144.
[6] M.D. Demetriou, M.E. Launey, G. Garrett, J.P. Schramm, D.C. Hofmann, W.L. Johnson, R.O. Ritchie, Nat. Mater. 10 (2011) 123–128. [7] Q. He, Y.Q. Cheng, E. Ma, J. Xu, Acta Mater. 59 (2011) 202–215. [8] M.E. Launey, D.C. Hofmann, W.L. Johnson, R.O. Ritchie, Proc. Natl. Acad. Sci. USA 106 (2009) 4986–4991. [9] D.C. Qiao, G.J. Fan, P.K. Liaw, H. Choo, Int. J. Fatigue 29 (2007) 2149–2154. [10] Y. Nakai, S. Hosomi, Mater. Trans. JIM 48 (2007) 1770–1773. [11] L. Huang, D.C. Qiao, B.A. Green, P.K. Liaw, J.F. Wang, S.J. Pang, T. Zhang, Intermetallics 17 (2009) 195–199. [12] U.K. Mudali, S. Baunack, J. Eckert, L. Schultz, A. Gebert, J. Alloys Comp. 377 (2004) 290–297. [13] K. Asami, C.L. Qin, T. Zhang, A. Inoue, Mater. Sci. Eng. A 375–377 (2004) 235– 239. [14] C.L. Qin, K. Asami, T. Zhang, W. Zhang, A. Inoue, Mater. Trans. JIM 44 (2003) 749–753. [15] V.R. Raju, U. Kühn, U. Wolff, F. Schneider, J. Eckert, R. Reiche, A. Gebert, Mater. Lett. 57 (2002) 173–177. [16] W.H. Peter, R.A. Buchanan, C.T. Liu, P.K. Liaw, M.L. Morrison, J.A. Horton, C.A. Carmichael Jr., J.L. Wright, Intermetallics 10 (2002) 1157–1162. [17] X.H. Chen, B.Y. Zhang, G.L. Chen, X.D. Hui, X.R. Xing, Intermetallics 19 (2011) 1913–1918. [18] P. Yu, K.C. Chan, L. Xia, H.B. Yu, H.Y. Bai, Mater. Trans. JIM 50 (2009) 2451– 2454. [19] Y. Zhang, X.H. Chen, G.L. Chen, X.C. Zhang, Z.D. Wang, China patent ZL200710120355.4, 2009. [20] X.H. Chen, B.Y. Zhang, G.L. Chen, Y. Zhang, X.D. Hui, Z.P. Lu, X.J. Liu, Y. Xu, X.R. Xing, Intermetallics 18 (2010) 2034–2038. [21] X.H. Chen, Y. Zhang, X.C. Zhang, X.D. Hui, B.C. Wei, G.L. Chen, Electrochem. Solid-State Lett. 10 (2007) 21–23. [22] M.L. Wang, G.L. Chen, X. Hui, Y. Zhang, Z.Y. Bai, Intermetallics 15 (2007) 1309– 1315. [23] A. Peker, W.L. Johnson, Appl. Phys. Lett. 63 (1993) 2342–2344. [24] K.L. Choy, J.F. Durodola, B. Derby, C. Ruiz, Composites 26 (1995) 531–539. [25] R.D. Conner, R.B. Dandliker, W.L. Johnson, Acta Mater. 46 (1998) 6089–6102. [26] Y. He, R.B. Schwarz, D.G. Mandrus, J. Mater. Res. 11 (1996) 1836–1841. [27] H.E. Boyer, T.L. Gall, Metals Handbook, Desk Edition., American Society for Metals, Metals Park, Ohio, 1985. [28] Y.S. Touloukian, R.K. Kirby, R.E. Taylor, P.D. Desai, Thermal Expansion: Metallic Elements and Alloys, IFI/Plenum, New York, 1975. [29] J. Zarzycki, Glasses and Vitreous State, Cambridge University Press, Cambridge, 1991.