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Construction of metallurgical interface with high strength between immiscible Cu and Nb by direct bonding method
Journal of Alloys and Compounds 723 (2017) 1053e1061
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Construction of metallurgical interface with high strength between immiscible Cu and Nb by direct bonding method Xinchang Pan, Jie Zhang, Yuan Huang*, Yongchang Liu Institute of Advanced Metallic Materials, School of Material Science and Engineering, Tianjin University, Tianjin, 300072, PR China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 27 March 2017 Received in revised form 6 June 2017 Accepted 28 June 2017 Available online 30 June 2017
Cu/Nb joints are promising to be used in resistance to radiation damage because of their outstanding properties. Due to the immiscibility in the Cu-Nb system, metallurgical interface is difficult to construct directly between Cu and Nb. In this paper, a direct bonding method was used to construct the metallurgical interface directly between Cu and Nb rods and prepare Cu/Nb joints. The key point of the method is to anneal Cu-Nb assemblies at the temperature between 92% and 98% of the melting point of Cu (TmCu ¼ 1356 K) for 3 h under a pressure of 106 MPa. The maximum tensile strength and bending strength of the obtained Cu/Nb joints are 222 MPa and 47 MPa, respectively. The Micro-test results show that a diffusion layer with a thickness of about 36 nm forms between Cu and Nb, indicating that a metallurgical interface is successfully constructed between Cu and Nb. The high temperature structure of Cu at the temperature close to the TmCu provides diffusion paths probably. In addition, a thermodynamic model was established for the direct bonding of Cu and Nb. Through the thermodynamic calculations and differential scanning calorimetry (DSC) tests, the storage energies in the Cu and Nb rods are proved to serve as the thermodynamic driving force for the diffusion between Cu and Nb. © 2017 Elsevier B.V. All rights reserved.
Keywords: Immiscible Cu-Nb system Metallurgical interface Direct bonding Diffusion Thermodynamic mechanism
1. Introduction The industrial components based on the immiscible Cu-Nb system have strong practical application values due to their outstanding comprehensive performances such as resistance to radiation damages, high strength, good ductility, high thermal stability and conductivity simultaneously [1e3]. For example, Cu/ Nb joints can be used in the thermonuclear fusion reactors, where Nb can be used as the plasma facing material (PFM) to resist radiation damage, and Cu can be used as the hot sink to dissipate the heats generated in the reactor. At present, the Cu/Nb joint with high strength is still an unsettling question because of the immiscibility in the Cu-Nb system. According to the present data, the solubility of Nb in Cu is only 0.096 wt% at 293 K [4]. Some methods for joining two different immiscible metals have been reported, which mainly include the diffusion welding process [5e8], the hot-rolling process and mechanical bonding [9]. For example, the diffusion welding have been employed to join the immiscible system such as Mo-Cu, W-Cu, W-Ag, Nb-Cu, Cu-Ag, Al-Be and Cu-Ta [7]. It need to be
* Corresponding author. E-mail address: [email protected] (Y. Huang). http://dx.doi.org/10.1016/j.jallcom.2017.06.314 0925-8388/© 2017 Elsevier B.V. All rights reserved.
pointed out that third metals are usually used as the middle interlayer [5e8,10,11] for the immiscible metals during the diffusion welding and the hot-rolling process. However, the methods using interlayer metals often undermined the consistency of the composition, bring some unnecessary performances. For example, Nickel produce ferromagnetism, and affect the operation of the equipments. The direct metallurgical bonding interfaces between the immiscible elements are not being realized now. This work is carried out with an aim to realize the high strength connection between the Cu and Nb rods without using interlayer metals by using a direct bonding method. The diagram of the direct bonding process of the Cu and Nb rods was shown in Fig. 1, during which the end surfaces of the Cu and Nb rods were polished firstly. Then the two rods were assembled coaxially according to the way that the polished end surfaces contact with each other. Finally, the assembled specimen was annealed under a certain pressure at a high temperature close to the melting point of Cu. We believe that the diffusion can be induced and the metallurgical bonding interface can be constructed directly between the immiscible metals through the above direct bonding process. The reasons can be described as follows. First, being different from the single crystal structure at the room temperature, some low melting point metals have a variety of crystal structure when being at the
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Fig. 1. Process of the connection of the Cu/Nb rods by direct bonding method.
temperature close to the melting point. For example, Ag have facecentered cubic (fcc), body-centered cubic (bcc), hexagonal closepacked (hcp) and amorphous structure at the solidification process [12], which can probably provide the diffusion paths for the refractory metals such as Mo, W, Nb. Secondly, in thermodynamics, the storage energies (Es) in the materials maybe released at the high temperature to overcome the positive heat of reaction and serve as the thermodynamic drive force for the diffusion. Finally, the pressure can also contribute to the diffusion. We think that the annealing temperature maybe the most crucial for the direct bonding method. To verify the above ideas, the joining of Cu and Nb was carried out by the direct bonding method at different temperatures. Then the microstructures of the Cu/Nb joints were characterized by a HITACHI S4800 field emission scanning electron microscope (SEM) and a Tecanai G2 F20 S-Twin transmission electron microscope (TEM) equipped with an Energy Dispersive X-Ray (EDX) Spectroscopy. Additionally, a new thermodynamic model was also constructed for the direct bonding method on basis of the Miedema theory. Combining differential scanning calorimetry (DSC) test, the model was used to verify that the storage energies (Es) can serve as the thermo dynamical driving force for the diffusion between Cu and Nb. The study verified the feasibility of the direct bonding method for the immiscible metals by experiment and theory finally. 2. Experimental procedures The base materials used for the direct bonding were commercially pure Cu rod (99.95 wt%) and commercially pure Nb rod (99.97 wt%), the sizes of which were all f 10 mm 25 mm. Before being connected, the end-surfaces of Cu and Nb rods were carefully ground flat by sandpaper down to 400, 600, 800, 1000, 1500 and 2000 grits and then polished by 0.5 mm Al2O3 powder, the aim of which was to remove the oxide films and the hardening films on the surfaces of the metal rods. The two rods were assembled coaxially in a fixture made of high temperature resistant materials as shown in Fig. 2, where the polished end-surfaces contacted with each other. Afterwards, the nuts of the fixture were tightened to the specified torque (10 N/m) with a torque wrench, and the Cu-Nb assembly was correspondingly applied to the pressure of 106 MPa at this time. The pressure shouldn't make significant plastic deformation for the specimens. It should be noted that a suitable pressure is conducive to the diffusion since it help to keep
Fig. 2. Schematic diagram of the fixture. 1-molybdenum plate; 2-quartz plate; 3-Cu rod; 4-Nb rod; 5-bolt; 6-nut; 7-pressure transducer.
the existence of the high temperature structure of Cu, and higher pressures may inhibit the rapid diffusion of atoms [13]. Then, the specimens were placed in the tube type annealing furnace (KTL-1600) with the fixture for annealing. The annealing temperatures were designed to be 1193 K, 1223 K, 1253 K, 1323 K, 1353 K and 1373 K. The protecting atmosphere was hydrogen. Fig. 3
Fig. 3. Schematic diagram of heat treatment process.
X. Pan et al. / Journal of Alloys and Compounds 723 (2017) 1053e1061
showed the diagram of heat treatment process. When the annealing finished, the Cu/Nb joints were obtained finally. Fig. 4 showed the obtained Cu/Nb joint. The tensile strength and bending strength of the joints were tested using WDW-20 universal machine at a velocity of 5 105 m/s. The fracture morphologies of tensile specimens were examined using the SEM, which is equipped with an Oxford energy dispersive spectrometer (EDS). The microstructure of the asconstructed Cu/Nb interface was observed by the TEM, where the test specimens were prepared by the method of ion-beam thinning. Differential scanning calorimetry (DSC3þ) was employed to test the storage energies (Es) in the Cu and Nb rods. Firstly, some of the Cu and Nb rods were pre-annealed at suitable temperatures to release the possible storage energy, while the rest are not carried out annealing. Then pre-annealed specimens and un-annealed specimens were all tested by using the DSC. Finally, the storage energies in the two metals could be concluded by the comparison of all the obtained DSC curves. The determination basis for the storage energy is that exothermic peaks shouldn't appear in the DSC curves for pre-annealed specimens since the storage energies have been released, and exothermic peaks should appear in the DSC curves for un-annealed specimens. 3. Results and discussion 3.1. HRTEM analysis of the as-constructed Cu/Nb interface Fig. 5 (a)-(b) is the HRTEM images of the cross-section of the Cu/ Nb joint obtained at 1253 K for 3 h. The insets of Fig. 5 (a) are the corresponding selected area electron diffraction (SAED) patterns of the bright area and the dark area in Fig. 5 (a) respectively. The indexing results of the patterns reveal that the bright area is Cu and the dark area is Nb. Fig. 5 (b) is the partially enlarged view of the region marked with white rectangular frame in Fig. 5 (a). It shows the high resolution lattice image of the Cu/Nb interface, from which it can be seen that a crystal interface has formed between Cu and Nb. Fig. 5 (d) presents the EDX line-scanning compositional profile of the Cu/Nb interface along the red line marked in Fig. 5 (c). According to Fig. 5 (d), it can be concluded that a mutual diffusion has occurred between Cu and Nb and the thickness of the diffusion layer is about 36 nm, which is a qualitative leap for immiscible metal systems. The result suggests that a metallurgical bonding interface has been successfully constructed between Cu and Nb during the direct bonding process. 3.2. Tensile strength test of Cu/Nb joints During the tensile test, it is found that the Cu/Nb joints obtained at 1193 K, 1223 K, 1353 K and 1373 K have almost negligible tensile strength as compared with the Cu/Nb joints obtained at 1253 K and
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1323 K. The load-displacement curves of tensile tests for the Cu/Nb joints obtained at 1253 K and 1323 K are shown in Fig. 6 (a) and (b), respectively. The effective fracture area (S) is measured with the software of Image-Pro Plus. The tensile fracture strength (s) is calculated according to equation (1):
s ¼ F=S
(1)
where, F is the maximum tensile load. All the results of the Cu/Nb joints including the maximum tensile load, the effective fracture area and the tensile strength are listed in Table 1. As shown in Table 1, the average tensile fracture strength of the Cu/Nb joints obtained at 1253 K and 1323 K are about 184 MPa and 193 MPa, respectively. The maximum tensile strength of about 222 MPa is obtained in the Cu/Nb joint bonded at 1253 K. These tensile strengths are close to or higher than that of the pure Cu (200e300 MPa) [14]. Additionally, the average maximum tensile load and the effective fracture area of the Cu/Nb joints obtained at 1253 K are obviously larger than those of the joints obtained at 1323 K, which means that the load capability of the Cu/Nb joints obtained at 1253 K is higher than that of the joints obtained at 1323 K. Considering the results that Cu and Nb rod can't be joined at 1193 K, 1223 K, 1353 K and 1373 K, it can be concluded that the annealing temperatures between 1253 K and 1323 K are the effective temperature for the direct bonding of the Cu and Nb rods. If the annealing temperature is outside the effective temperature range 1253e1323 K, it is difficult to connect the Cu and Nb rods. Obviously, the direct bonding method of the immiscible metals is very sensitive to the annealing temperatures which need to be controlled accurately. Since the annealing temperatures are close to the melting point of Cu (TmCu) and considering that it is impossible to connect the Cu and Nb rods when the annealing temperatures are higher than TmCu (1356 K), TmCu is used as the standard to evaluate the effective temperature range for the Cu/Nb bonding system. The specific method is to calculate the ratio of annealing temperature to TmCu. So, the effective temperature range 1253e1323 K can be expressed as 0.92 TmCu-0.98 TmCu, which become a very intuitive concept. The above results can be supported by other references. For example, He and his coworkers [15] annealed an immiscible Mo/Zn bilayered sample at 673 K for 5 h under the protection of Ar gas. Since the melting temperature of Zn, TmZn, is 692.53 K, the annealing temperature of 673 K is 0.97 TmZn, which is in the effective temperature range presented in this paper. The EDX linescan profiles performed in the annealed sample show that the difficulty for the element diffusion of immiscible alloy systems is overcome and Mo atoms diffuse a long distance (~600 nm) in Zn layer at the annealing temperature.
Fig. 4. Cu/Nb joint obtained through the direct bonding method.
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Fig. 5. HRTEM analysis of the cross-section of the Cu/Nb joint bonded at 1253 K: (a) HRTEM image of the Cu/Nb interface; (b) the partially enlarged view of the region marked with white frame in (a); (c) drift-corrected spectrum image scanning; (d) the curve of diffusion depth along the red line marked in (c). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
Fig. 6. Load-Displace curves of the tensile tests of the Cu/Nb joints obtained at (a) 1253 K and (b) 1323 K, respectively.
Table 1 Tensile results of the Cu/Nb joints bonded at different temperatures. Sample Annealing temperature (K)
Maximum Effective fracture Tensile fracture tensile load (N) area (mm2) strength (MPa)
1 1253 2 1253 3 1253 Average
4507 4186 5241 4645
20.24 23.18 33.89 25.77
222.67 180.59 154.65 185.97
4 1323 5 1323 6 1323 Average
2266 3055 3025 2782
10.98 15.67 16.90 14.52
206.38 194.96 178.99 193.44
3.3. Bending strength test of Cu/Nb joints The bending strengths of the Cu/Nb joints are obtained by threepoint bending test in this study. The tests were carried out on three samples which were prepared through the same direct bonding process, where the annealing temperature was 1253 K, the bonding time was 3 h and the bonding pressure was 106 MPa. The load-displacement curves of bend tests are shown in Fig. 7 and the corresponding results are listed in Table 2. As shown in Table 2, the average bending load and bending strength of the Cu/ Nb joints can reach 836 N and 42 MPa, respectively. The above mechanical properties of the Cu/Nb joints also prove that the as-constructed Cu/Nb interfaces have high strength.
X. Pan et al. / Journal of Alloys and Compounds 723 (2017) 1053e1061
Fig. 7. Load-displacement curves of the bending tests for the Cu/Nb joints obtained at 1253 K.
Table 2 Results of the bending tests for the Cu/Nb joints obtained at 1253 K. Sample Temperature (K)
Max load (N)
Fracture area (mm2)
Max strength (MPa)
1 1253 2 1253 3 1253 Average
803.7 873.2 832.6 836.5
18.15 18.31 22.69 19.72
44.28 47.69 36.69 42.89
3.4. SEM analysis of the tensile fractures of Cu/Nb joints The SEM images shown in Fig. 8 are the tensile fractures of three Cu/Nb joint specimens annealed similarly at 1253 K for 3 h under the pressure of 106 MPa. In Fig. 8, many equiaxed dimples with different sizes and micro-holes can be obviously observed. These results indicate that the fractures of the Cu/Nb joints obtained through the direct bonding method are mainly plastic when being stretched. Fig. 8 (a)-(c) are the fractures on end-surfaces of the Cu rods of the Cu/Nb joints, where many deep dimples can be seen. Fig. 8 (d)(f) are the fractures on end-surfaces of the Nb rods of the Cu/Nb joints, where many shallow dimples can be seen. Since dimple is a peculiar feature of the fracture morphology of pure Cu, these
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results mean that the fractures occur in a part of the Cu rod near to the Cu/Nb interface, indicating that the bonding strength of the Cu/ Nb metallurgical interface is very high. In addition, at the area marked with white rectangular frame in Fig. 8 (d), a steady river pattern can be seen, which is characteristic of brittle cleavage fracture in morphology. Since the cleavage fracture usually occurs in the bcc lattice metal and Nb is just bcc lattice metal, it can be concluded that the Nb rod is fractured during the tensile at some time, which also means that the bonding strength of the Cu/Nb metallurgical interface is very high. The compositions of the dimples of the specimens have been analyzed by means of the EDX spectrometer in the SEM, which are shown in Fig. 9. Fig. 9 (a) is the fracture on the end-surface of the Nb rod of the Cu/Nb joint. Fig. 9 (b) presents the EDX line-scanning composition profile along the white line marked in Fig. 9 (a). It can be seen from Fig. 9 (b) that the two composition curves show periodic fluctuation, which are caused by the structure of the dimples. At the bottoms of the dimples, the concentration of Nb is higher than that of Cu, conversely, the concentration of Cu is higher than that of Nb at the tops of the dimples. Considering the morphology of the dimples still remain the peculiar feature of pure Cu, it can be concluded that the fracture of the Cu/Nb joints mainly occur at the end of the Cu rod. The coexistence of Nb and Cu in the dimple prove once again the mutual diffusion occurring between the Cu and Nb rods during the bonding process. 3.5. Diffusion mechanism of the direct bonding method As discussed above, the temperatures between 1253 K and 1323 K, namely 0.92 TmCu-0.98 TmCu are the effective temperature for the direct bonding of the Cu and Nb rods. Only in that effective temperature range can the mutual diffusion occurs between the Cu and Nb rods and form the metallurgical bonding interface. The reason can be described as following. In the temperature 0.92 TmCu-0.98 TmCu, Cu is in semi-solid state and the Cu may have phase transition [16,17]. In this state, 60% of Cu can be solid phase and 40% of Cu is in a molten-state [16]. Molecular dynamics simulations (MD) [18] proves that the face-centered cubic (fcc), body-centered cubic (bcc), hexagonal close-packed (hcp) structures and amorphous phase can coexist in Cu in semi-solid state [19]. And the simulation result also shows that during the processes of solidification of liquid Cu and amorphous crystallization, the crystal embryo of bcc is always formed firstly [20]. In addition, the lattice constant of Cu at 1253 K is about 0.369 nm,
Fig. 8. SEM images of tensile fracture surface corresponding to the Cu and Nb rods of three Cu/Nb joints: (a) and (d) specimen 1, (b) and (e) specimen 2 and (c) and (f) specimen 3, respectively.
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Fig. 9. Analysis of the composition of the dimples in the Nb rod of the Cu/Nb joint. (a) SEM of fracture morphology and (b) corresponding line-scanning composition profile along the white line marked in (a).
which is larger than that of Cu room temperature (0.361 nm). Correspondingly, the atomic gap of Cu will also become larger at 1253 K [21]. According to the above discussion, there may be several kinds of diffusion paths for the diffusions of Cu and Nb. Firstly, bcc structure of Cu in semi-solid state and larger atomic gap at high temperature can serve as diffusion path for the Nb atoms. Secondly, there will be some spot-defects in the Cu melt, such as vacancies [22,23]. At this time, the Nb atoms may occupy the vacancies in the Cu melt by thermal motion, reaching the purpose of diffusion. Finally, since the crystal structure of Nb is bcc, the bcc structure appears in the Cu in semi-solid state may also caused some Cu atoms diffuse into Nb. But, in general, the diffusion of Nb into Cu should be dominant because the Cu in semi-solid state can provide more diffusion paths. If the temperature is lower than 0.92 TmCu or higher than 0.98 TmCu, the diffusion path may be inadequate. The specific diffusion mechanism needs to be further explored. 3.6. Thermodynamic mechanism for the direct bonding of Cu and Nb 1. Thermodynamic model In this work, a thermodynamic model is constructed for the direct bonding method used for the connection of the immiscible Cu and Nb metals. According to the process as shown in Fig. 1, the total energy change during the bonding process can be expressed as follows:
DGtotal ¼ DGðCu=NbÞ þ DGint þ DGbulkðCu;NbÞ þ Es þ P$DV
(2)
where,
DGtotal is the total Gibbs free energy during the direct bonding process;
and DGint . On the other hand, DGðCu=NbÞ and DGint are really the
Gibbs free energy of the alloying of Cu and Nb, namely DGalloying. So, equation (2) can be rewritten as following:
DGinitial ¼ Es þ P$DV þ DGbulkðCu;NbÞ DGalloying ¼ DGðCu=NbÞ þ DGint
tween Cu and Nb, and can be described as follows: in NbÞ
Among which, DGðCu
þ DGðNb
in NbÞ
in CuÞ
(3)
is the Gibbs free energy caused by
the diffusion of Cu into Nb, and DGðNb in CuÞ is the free energy caused by the diffusion of Nb into Cu. For the solid-solution alloy system, these items are usually negative, the corresponding diffusion can occur automatically. However, for the immiscible alloy system, the items are usually positive, which need to get energy support from outside; DGint is the interfacial free energy between Cu and Nb, which
(4)
Equation (4) is the thermodynamic model constructed for the direct bonding of Cu and Nb. If DGinitial is larger than DGalloying , the diffusion and alloying between Cu an Nb can occur and proceed. In other words, the initial free energy, DGinitial can serves as the thermodynamic driving force for the diffusion and alloying at this time. The calculation formulas for the above thermodynamic quantities are based on Miedema model and Alonso method [24e26]. Generally, the values of DGðCu=NbÞ are calculated by equations (3), (5) and (6).
DGðCu
in NbÞ
¼ DH ðCu
in NbÞ
T 0 DS
(5)
DGðNb
in CuÞ
¼ DH ðNb
in CuÞ
T 0 DS
(6)
Where, DH ðCu in NbÞ /ðDH ðNb in CuÞ Þ and DS are the enthalpy and entropy of formation, respectively.
DS ¼ RðX A lnX A þ X B lnX B Þ
DGðCu=NbÞ is the Gibbs free energy caused by the diffusion be-
DGðCu=NbÞ ¼ DGðCu
also need to get energy support from outside; DGbulkðCu;NbÞ is the Gibbs free energy of the mechanical mixture of metal Cu and Nb; P$DV is the energy caused by the pressure; E s are the storage energies in the Cu and Nb rods, which could be measured by the DSC test. According to the statements above, Es , P$DV and DGbulkðCu;NbÞ form the initial Gibbs free energy of the Cu-Nb bonding system, namely DGinitial, which can provide energy support for DGðCu=NbÞ
(7)
where, R is the gas constant; XA and XB are the atomic concentrations of metal A (Cu) and B (Nb) in the alloy, respectively. According to Miedema et al., the enthalpy of formation can be calculated as follows [27e29]:
DH A
in BðB in AÞ
¼ DH cA
in BðB in AÞ
þ DH e þ DH s
(8)
where, DH c , DH e and DH s correspond to chemical, elastic and structural contributions, respectively. Since the calculated value of DH e is 6.06 104 kJ/mol in this paper and the contribution of DH s to the formation enthalpy are very small [30], the elastic and structural terms are ignored. The chemical contribution, DH c , is
X. Pan et al. / Journal of Alloys and Compounds 723 (2017) 1053e1061
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described as follows [31]: A in BðB in AÞ
DH cA
A in B
in B
¼ xA fAB DH sol
The solution heat of DH Cu sol Ref. [31]. 2=3
DH Asolin
B
¼
2PV A
h
gchem ¼ ss
(9) in Nb
and DH Nb sol
in Cu
are given by
2. DSC test of storage energies
(10)
The values of Es (the storage energy) in equation (4) have been measured by the DSC test according to the process described in Section 2. Before the DSC test, the Cu rod was cut into two rods with the same quality, and named as specimen 1 and specimen 2, respectively. The same process was implemented in the Nb rod, and the obtained two rods were named as specimen 3 and specimen 4. Specimen 2 and 4 were annealed at 1323 K for 6 h and 1873 K for 6 h to release the possible storage energy before the DSC tests. The final DSC results are shown in Fig. 10, where curve 1, 2 3 and 4 corresponding to specimen 1, 2, 3 and 4, respectively. It can be seen from Fig. 10 that the apparent exothermic peaks appear at curve 1 and curve 3, and the areas of the peaks are 10.36 kJ/mol and 49.30 kJ/mol. Conversely, there are no obvious exothermic peaks appearing in curve 2 and 4. Comparing the curves, it can be concluded the storage energies in the Cu and Nb rods are 10.36 kJ/ mol and 49.30 kJ/mol, respectively. It also can be seen from Fig. 10 that the storage energy in the Cu rod is released in the temperature range of 873e1073 K and the storage energy in the Nb rod is released in the temperature range of 1073e1273 K. So, when the annealing temperature is higher than 1273 K, the total released storage energy in the Cu-Nb bonding system should be the sum of the storage energies of the Cu and Nb rods.
where, P, Q and R are empirical constant; 4A and 4B are the chemical potentials of metal A and B; nwsA and nwsB are the Wigner-Seitz concentration of the cell boundary of metal A and B, respectively. The f BA is a function which accounts for the degree to which atoms of type A are surrounded by atoms of type B, given by Ref. [32].
(11)
where, s is short-range order parameter of 5 for Cu-Nb system and C SA /C SB is the surface concentration of A/B atoms. The interfacial free energy, DGint is calculated by Ref. [32].
DGint ¼ aA DGint;A þ aB DGint;B
(12)
where, aA and aB are the proportions of the interface atoms in the alloy of A and B, respectively. In our calculation, aA and aB are deduced [32].
aA ¼
PnA
DdA;i
xA Pi¼1 NA i¼1
dA;i
PnB
; aB ¼ xB Pi¼1 NB
DdB;i
i¼1
3. Thermodynamic driving forces for the diffusion and interface alloying
(13)
dB;i
Since the Cu-Nb system is immiscible, the DGbulkðCu;NbÞ term in the thermodynamic model described by equation (4) can be considered as zero. Additionally, according to Hooke's law, the volume change made by the pressure (P) in the Cu/Nb contacting surface, namely DV, is 106 m3. So, the value of the P$DV term in equation (4) is 0.32 kJ/mol. The other thermodynamic quantities including DGðCu in NbÞ , DGðNb in CuÞ and DGint are calculated ac-
The interfacial atomic free energy of 1 mol of metal is given by Ref. [32].
DGint;i ¼ S f ;i *gAB ss
(14)
s;0 s;0 þ gchem gAB ss ¼ 0:15 gA þ gB ss
(16)
2=3
C0V A
where, C 0 is a constant (4:5 108 ), and V A is the atomic volume of component A in the alloy. All the parameters used in the calculations are listed in Table 3 and Table 4.
i 1=3 1=3 2 ð4A 4B Þ2 þ QP nwsA nwsB R P 1=3 1=3 nwsA þ nwsB
2 2 C SB f BA ¼ C SA 1 þ s C SA
DH sol
(15)
cording to equation (5)e(16). Fig. 11 is the final calculated Gibbs free energy diagrams of the Cu-Nb bonding system according to the as-constructed thermodynamic model. According to Fig. 11, the green line is the alloying free energy
Where, gs;0 and gs;0 B are the surface energy of metals A and B, A being 1825 and 2700 mJ/m2 for Cu and Nb, respectively; gchem is ss given by Ref. [32].
Table 3 Calculation parameters of DGðCu=NbÞ [26]. e means the data did not exist. 2=3
metal/system
Vi
Cu Nb Cu-Nb
3.70 4.89 e
(cm2)
1=3
nws (d.u.)1/3
F (V)
Q/P
P
R (J/mol*K)
s
T0 (K)
1.47 1.68 e
4.55 4.72 e
e e 9.4
e e 14.1
e e 8.314
e e 5
e e 300
di (104cm)
Ddi (108cm)
Sf (105m2)
C0 (108)
5 11 e
e e 5
1.66 2.20 e
e e 4.5
Table 4 Calculation parameters of DGint [26]. e means the data did not exist. 2=3
metal/system
Vi
Cu Nb Cu-Nb
3.70 4.89 e
(cm2)
gs;0 (mJ*m2) i
DH Cu sol
1825 2700 e
e e 31.56
in Nb
(kJ/mol)
DH Nb sol e e 40.22
in Cu
(kJ/mol)
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the Cu and Nb rods completely, the diffusion and alloying in the CuNb bonding system can occur and proceed. It need to be pointed out that only a large thermodynamic driving force is not enough for the diffusion between Cu and Nb, sufficient diffusion paths are also necessary. That is the reason why the joining of the Cu and Nb rods by the direct bonding method should be carried out in a narrow and sensitive temperature range of 0.92 TmCu-0.98 TmCu. 4. Conclusions
Fig. 10. DSC curves of Cu and Nb rods: 1#-non-annealed Cu rod; 2#-annealed Cu rod; 3#-non-annealed Nb rod; 4#- annealed Nb rod.
Fig. 11. Calculated diagram of the initial energy and alloying energy of the Cu-Nb bonding system.
(DGalloying ) of the Cu-Nb bonding system and the maximum value is 19.48 kJ/mol. The red line is the initial free energy (DGinitial ) of the Cu-Nb bonding system at 873e1073 K, the maximum value of which is 10.36 kJ/mol. It can be seen that the DGinitial is much lower than the DGalloying in the most of the composition of Cu, indicating that the bonding system does not have enough thermodynamic driving force to ensure that the diffusion and alloying occur between the Cu and Nb rods. The main cause resulting in the low initial free energy is that only the storage energy in the Cu rod is released at this time. The blue line in Fig. 11 is the initial free energy (DGinitial ) of the Cu-Nb bonding system at the temperature higher than 1073 K. It can be seen that the DGinitial term is much higher than the DGalloying term, which is caused by the co-release of the storage energies of the Nb and Cu rods. The pink arrow in Fig. 11 shows the minimum difference of 40.51 kJ/mol between DGinitial and DGalloying , which indicates that the bonding system has an enough thermodynamic driving force to ensure that the diffusion and alloying occur at the Cu/Nb interface at the temperature higher than 1073 K. Since the calculated P$DV is very small and the DGbulkðCu;NbÞ term is zero, the initial energy, DGinitial depends almost entirely on the release of the storage energies (Es ) in Cu and Nb rods. In other words, Es serves as the thermodynamic driving force for the diffusion and alloying occurring at the Cu/Nb interface during the direct bonding process. Only the Cu-Nb bonding system is annealed at sufficiently high temperature to release the storage energies in
In this paper, a direct bonding method was used to construct the metallurgical interface directly between immiscible Cu and Nb without using an interlayer metal. The construction processes include the polish for the end-surfaces of the Cu and Nb rods, the coaxial assembly of the Cu and Nb rods, and sequential annealing at high temperature for 3 h under a pressure of 106 MPa. Through the method, a Cu/Nb metallurgical interface was successfully constructed and the corresponding Cu/Nb joint was obtained. The maximum tensile strength and bending strength of the obtained Cu/Nb joints are about 222 MPa and 47 MPa, respectively, which have reached a very high level. The micro-test results show that thickness of the Cu/Nb diffusion layer is about 36 nm. Additionally, a thermodynamic model was established for the direct bonding of the immiscible metals by using the Miedema theory. Through the calculations based on the model and differential scanning calorimetry (DSC) tests, the storage energies of Cu and Nb rods released at high temperature are confirmed to be able to overcome the positive reaction heat of Cu-Nb system, and serve as the thermodynamic driving force for the diffusion and alloying occurring at the Cu/Nb interface. Finally, it need to be pointed out that the key point of the direct bonding method is the annealing temperature, which should be close to the melting point of Cu (TmCu ¼ 1356 K). On the one hand, the high-temperature structure of Cu at the temperature close to the TmCu can provide the diffusion paths for Cu and Nb probably. On the other hand, the storage energies of Cu and Nb rods can be fully released to drive the diffusion. The experimental results show that the suitable annealing temperature should be between 0.92 TmCu0.98 TmCu, which is very sensitive and need to be controlled accurately. Acknowledgement This work was supported by the National Natural Science Foundation of China (51171118 and 51471114), and the Science and Technology Support Project of Tianjin City (11ZCKFGX03800). References [1] X. Zhang, N. Li, O. Anderoglu, H. Wang, Nuclear instruments and methods in physics research section B: beam interactions with materials and atoms, Nucl. Instrum. Methods Phys. B 261 (2007) 1119. [2] A. Misra, M.J. Denkowicz, The radiation damage tolerance of ultra-high strength nanolayered composites, JOM 59 (2007) 62e65. [3] W.Z. Han, A. Misra, N.A. Mara, Role of interfaces in shock-induced plasticity in Cu/Nb nanolaminates, Philos. Mag. 91 (2011) 4172. [4] PonoB.И.A., ШИPЯCBa H.B. // ЖypHaл HeopraHHqeckoй Xиmии. T.6. 10. (1961) 2334e2340. [5] W. Elthallabawy, T. Khan, Liquid phase bonding of 316L stainless steel to AZ31 magnesium alloy, Mater. Sci. Technol. 27 (2011) 22e28. [6] M. Ding, S.S. Liu, Y.C. Wang, TIG-MIG hybrid welding of ferritic stainless steels and magnesium alloys with Cu interlayer of different thickness, Mater. Des. 88 (2015) 375e383. [7] F.M. Zhou, F.Q. Zhang, F.Y. Song, Vacuum diffusion bonding and interfacial structure in Ta/Cu couple, Rare Metal Mater. Eng. 42 (2013) 1785e1789. [8] S. Wang, Y.H. Ling, J.J. Wang, Microstructure and mechanical properties of W/ Cu vacuum diffusion bonding joints using amorphous Fe-W alloy as interlayer, Vacuum 114 (2015) 58e65.
X. Pan et al. / Journal of Alloys and Compounds 723 (2017) 1053e1061 [9] J. Lin, H.F. Zhao, Mechanical and metallurgical bonding in friction stir welding joint of steel-aluminum dissimilar materials, J. Mech. Eng. 16 (2015) 106e112. [10] J.L. Fan, S.Z. Yang, T. Liu, Q. Zhou, Diffusion bonding process between W-Cu composite and Cu, Mater. Sci. Eng. Powder Metall. 2 (2015) 182e205. [11] J. Zhang, Q. Shen, G.Q. Luo, et al., Microstructure and bonding strength of diffusion welding of Mo/Cu joints with Ni interlayer, Mater. Des. 39 (2011) 81e86. [12] W. Huang, G.Y. Liang, In suit transformation of atomic arrangement of amorphous Ag during early crystallization: a molecular dynamic study[EB/ OL]. http://www.paper.edu.cn/releasepaper/content, 2012-01-21. [13] Z.H. Yan, H. Chen, X.M. Huang, Solid state amorphous reaction under high pressure, Chin. J. High Press. Phys. (1990), 4e1. [14] GB-T 4423-2007, Copper and Copper-alloy Cold-drawn Rod and Bar, Standards Press of China, 2007. http://www.spsp.gov.cn/page/P400/694.shtml. [15] Z.X. He, Y. Huang, F. He, Preparation of nanoporous molybdenum film by dealloying an immiscible MoeZn system for hydrogen evolution reaction, RSC Adv. 06 (2016) 15390e15393. [16] Z.Y. Du, Fundamentals of Welding Science: Fundamentals of Material Science, China Machine Press, 2012. [17] X.H. Yi, Properties of formation and growth of cluster structure during solidification processes of liquid metal Cu, Mater. Rev. B 29 (12) (2015) 122e127. [18] X.P. Li, Q.Y. Han, Phase structure transition of liquid copper during rapid cooling by MD simulation, Acta Metall. Sin. 31 (8) (1995) 356e362. [19] X.H. Yi, R.S. Liu, et al., Simulation study of effect of cooling rate on evolution of microstructures during solidification of liquid metal Cu, Acta Phys. Sin. 10 (2006) 5386e5392. [20] S. Alexander, T. McTague, All crystal Be bcc landau theory of solidification and
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crystal nucleation, Phys. Rev. Lett. 41 (10) (1978) 702e705. [21] H.L. Wang, X.X. Wang, Molecular dynamics simulation and analysis of bulk and surface melting processes for metal Cu, Acta Metall. Sin. 06 (2005) 568e572. [22] J.R. Cahoon, A modified “hole” theory for solute impurity diffusion in liquid metals, Metall. Mater Trans. A 28 (1997) 583e593. [23] C.N. Zhu, Y.L. Geng, B. Zhang, Research Progress on Diffusion of Metal Melt, vol. 06, Science China Press, 2012, pp. 619e630. [24] C.F. Wu, H.G. Li, Thermodynamic model of the binary alloys-miedema model, Shanghai Met. 33 (4) (2011) 1e5. [25] F.R. de Boer, R. Boom, A.R. Miedema, A.K. Niessen, Cohesion in Metals: Transition Metal Alloys, North-Holland, Amsterdam, 1989. [26] L. Viots, A.V. Ruban, H.L. Skriver, et al., The surface energy of metals, Surf. Sci. 411 (1998) 186e202. [27] A.K. Niessen, A.R. Miedema, Ber Bunscngcs, The enthalpy of solution for solid binary alloys of two 4d-transition metals, Phys. Chem. 7 (1983) 27e36. [28] J.M. Lopez, J.A. Alonso, Semi empirical theory of solid solubility in transition metal alloys, Z. Naturforsch. A 40 (1985) 199. [29] A.W. Weeber, The preparation of amorphous Ni-Zr alloys by different ballmilling techniques, Physica 45 (3) (1987) 349e352. [30] B.X. Liu, W.S. Lai, Z.J. Zhang, Solid-state crystal to amorphous transition in metal-metal multilayers and its thermodynamic and atomistic modelling, Adv. Phys. 04 (2001) 367e429. [31] A.K. Niessen, A.R. Miedema, F.R. de Boer, R. Boom, Enthalpies of formation of liquid and solid binary alloys based on 3d metals:ІІІ[J].Alloys of iron, Phys. B 151 (1983) 285e309. [32] V.P. Martinea, Mechanical alloying of Cu-Mo powder mixtures and thermodynamic study of solubility, Mater. Lett. 61 (4) (2007) 929e933.
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Construction of metallurgical interface with high strength between immiscible Cu and Nb by direct bonding method Xinchang Pan, Jie Zhang, Yuan Huang*, Yongchang Liu Institute of Advanced Metallic Materials, School of Material Science and Engineering, Tianjin University, Tianjin, 300072, PR China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 27 March 2017 Received in revised form 6 June 2017 Accepted 28 June 2017 Available online 30 June 2017
Cu/Nb joints are promising to be used in resistance to radiation damage because of their outstanding properties. Due to the immiscibility in the Cu-Nb system, metallurgical interface is difficult to construct directly between Cu and Nb. In this paper, a direct bonding method was used to construct the metallurgical interface directly between Cu and Nb rods and prepare Cu/Nb joints. The key point of the method is to anneal Cu-Nb assemblies at the temperature between 92% and 98% of the melting point of Cu (TmCu ¼ 1356 K) for 3 h under a pressure of 106 MPa. The maximum tensile strength and bending strength of the obtained Cu/Nb joints are 222 MPa and 47 MPa, respectively. The Micro-test results show that a diffusion layer with a thickness of about 36 nm forms between Cu and Nb, indicating that a metallurgical interface is successfully constructed between Cu and Nb. The high temperature structure of Cu at the temperature close to the TmCu provides diffusion paths probably. In addition, a thermodynamic model was established for the direct bonding of Cu and Nb. Through the thermodynamic calculations and differential scanning calorimetry (DSC) tests, the storage energies in the Cu and Nb rods are proved to serve as the thermodynamic driving force for the diffusion between Cu and Nb. © 2017 Elsevier B.V. All rights reserved.
Keywords: Immiscible Cu-Nb system Metallurgical interface Direct bonding Diffusion Thermodynamic mechanism
1. Introduction The industrial components based on the immiscible Cu-Nb system have strong practical application values due to their outstanding comprehensive performances such as resistance to radiation damages, high strength, good ductility, high thermal stability and conductivity simultaneously [1e3]. For example, Cu/ Nb joints can be used in the thermonuclear fusion reactors, where Nb can be used as the plasma facing material (PFM) to resist radiation damage, and Cu can be used as the hot sink to dissipate the heats generated in the reactor. At present, the Cu/Nb joint with high strength is still an unsettling question because of the immiscibility in the Cu-Nb system. According to the present data, the solubility of Nb in Cu is only 0.096 wt% at 293 K [4]. Some methods for joining two different immiscible metals have been reported, which mainly include the diffusion welding process [5e8], the hot-rolling process and mechanical bonding [9]. For example, the diffusion welding have been employed to join the immiscible system such as Mo-Cu, W-Cu, W-Ag, Nb-Cu, Cu-Ag, Al-Be and Cu-Ta [7]. It need to be
* Corresponding author. E-mail address: [email protected] (Y. Huang). http://dx.doi.org/10.1016/j.jallcom.2017.06.314 0925-8388/© 2017 Elsevier B.V. All rights reserved.
pointed out that third metals are usually used as the middle interlayer [5e8,10,11] for the immiscible metals during the diffusion welding and the hot-rolling process. However, the methods using interlayer metals often undermined the consistency of the composition, bring some unnecessary performances. For example, Nickel produce ferromagnetism, and affect the operation of the equipments. The direct metallurgical bonding interfaces between the immiscible elements are not being realized now. This work is carried out with an aim to realize the high strength connection between the Cu and Nb rods without using interlayer metals by using a direct bonding method. The diagram of the direct bonding process of the Cu and Nb rods was shown in Fig. 1, during which the end surfaces of the Cu and Nb rods were polished firstly. Then the two rods were assembled coaxially according to the way that the polished end surfaces contact with each other. Finally, the assembled specimen was annealed under a certain pressure at a high temperature close to the melting point of Cu. We believe that the diffusion can be induced and the metallurgical bonding interface can be constructed directly between the immiscible metals through the above direct bonding process. The reasons can be described as follows. First, being different from the single crystal structure at the room temperature, some low melting point metals have a variety of crystal structure when being at the
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Fig. 1. Process of the connection of the Cu/Nb rods by direct bonding method.
temperature close to the melting point. For example, Ag have facecentered cubic (fcc), body-centered cubic (bcc), hexagonal closepacked (hcp) and amorphous structure at the solidification process [12], which can probably provide the diffusion paths for the refractory metals such as Mo, W, Nb. Secondly, in thermodynamics, the storage energies (Es) in the materials maybe released at the high temperature to overcome the positive heat of reaction and serve as the thermodynamic drive force for the diffusion. Finally, the pressure can also contribute to the diffusion. We think that the annealing temperature maybe the most crucial for the direct bonding method. To verify the above ideas, the joining of Cu and Nb was carried out by the direct bonding method at different temperatures. Then the microstructures of the Cu/Nb joints were characterized by a HITACHI S4800 field emission scanning electron microscope (SEM) and a Tecanai G2 F20 S-Twin transmission electron microscope (TEM) equipped with an Energy Dispersive X-Ray (EDX) Spectroscopy. Additionally, a new thermodynamic model was also constructed for the direct bonding method on basis of the Miedema theory. Combining differential scanning calorimetry (DSC) test, the model was used to verify that the storage energies (Es) can serve as the thermo dynamical driving force for the diffusion between Cu and Nb. The study verified the feasibility of the direct bonding method for the immiscible metals by experiment and theory finally. 2. Experimental procedures The base materials used for the direct bonding were commercially pure Cu rod (99.95 wt%) and commercially pure Nb rod (99.97 wt%), the sizes of which were all f 10 mm 25 mm. Before being connected, the end-surfaces of Cu and Nb rods were carefully ground flat by sandpaper down to 400, 600, 800, 1000, 1500 and 2000 grits and then polished by 0.5 mm Al2O3 powder, the aim of which was to remove the oxide films and the hardening films on the surfaces of the metal rods. The two rods were assembled coaxially in a fixture made of high temperature resistant materials as shown in Fig. 2, where the polished end-surfaces contacted with each other. Afterwards, the nuts of the fixture were tightened to the specified torque (10 N/m) with a torque wrench, and the Cu-Nb assembly was correspondingly applied to the pressure of 106 MPa at this time. The pressure shouldn't make significant plastic deformation for the specimens. It should be noted that a suitable pressure is conducive to the diffusion since it help to keep
Fig. 2. Schematic diagram of the fixture. 1-molybdenum plate; 2-quartz plate; 3-Cu rod; 4-Nb rod; 5-bolt; 6-nut; 7-pressure transducer.
the existence of the high temperature structure of Cu, and higher pressures may inhibit the rapid diffusion of atoms [13]. Then, the specimens were placed in the tube type annealing furnace (KTL-1600) with the fixture for annealing. The annealing temperatures were designed to be 1193 K, 1223 K, 1253 K, 1323 K, 1353 K and 1373 K. The protecting atmosphere was hydrogen. Fig. 3
Fig. 3. Schematic diagram of heat treatment process.
X. Pan et al. / Journal of Alloys and Compounds 723 (2017) 1053e1061
showed the diagram of heat treatment process. When the annealing finished, the Cu/Nb joints were obtained finally. Fig. 4 showed the obtained Cu/Nb joint. The tensile strength and bending strength of the joints were tested using WDW-20 universal machine at a velocity of 5 105 m/s. The fracture morphologies of tensile specimens were examined using the SEM, which is equipped with an Oxford energy dispersive spectrometer (EDS). The microstructure of the asconstructed Cu/Nb interface was observed by the TEM, where the test specimens were prepared by the method of ion-beam thinning. Differential scanning calorimetry (DSC3þ) was employed to test the storage energies (Es) in the Cu and Nb rods. Firstly, some of the Cu and Nb rods were pre-annealed at suitable temperatures to release the possible storage energy, while the rest are not carried out annealing. Then pre-annealed specimens and un-annealed specimens were all tested by using the DSC. Finally, the storage energies in the two metals could be concluded by the comparison of all the obtained DSC curves. The determination basis for the storage energy is that exothermic peaks shouldn't appear in the DSC curves for pre-annealed specimens since the storage energies have been released, and exothermic peaks should appear in the DSC curves for un-annealed specimens. 3. Results and discussion 3.1. HRTEM analysis of the as-constructed Cu/Nb interface Fig. 5 (a)-(b) is the HRTEM images of the cross-section of the Cu/ Nb joint obtained at 1253 K for 3 h. The insets of Fig. 5 (a) are the corresponding selected area electron diffraction (SAED) patterns of the bright area and the dark area in Fig. 5 (a) respectively. The indexing results of the patterns reveal that the bright area is Cu and the dark area is Nb. Fig. 5 (b) is the partially enlarged view of the region marked with white rectangular frame in Fig. 5 (a). It shows the high resolution lattice image of the Cu/Nb interface, from which it can be seen that a crystal interface has formed between Cu and Nb. Fig. 5 (d) presents the EDX line-scanning compositional profile of the Cu/Nb interface along the red line marked in Fig. 5 (c). According to Fig. 5 (d), it can be concluded that a mutual diffusion has occurred between Cu and Nb and the thickness of the diffusion layer is about 36 nm, which is a qualitative leap for immiscible metal systems. The result suggests that a metallurgical bonding interface has been successfully constructed between Cu and Nb during the direct bonding process. 3.2. Tensile strength test of Cu/Nb joints During the tensile test, it is found that the Cu/Nb joints obtained at 1193 K, 1223 K, 1353 K and 1373 K have almost negligible tensile strength as compared with the Cu/Nb joints obtained at 1253 K and
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1323 K. The load-displacement curves of tensile tests for the Cu/Nb joints obtained at 1253 K and 1323 K are shown in Fig. 6 (a) and (b), respectively. The effective fracture area (S) is measured with the software of Image-Pro Plus. The tensile fracture strength (s) is calculated according to equation (1):
s ¼ F=S
(1)
where, F is the maximum tensile load. All the results of the Cu/Nb joints including the maximum tensile load, the effective fracture area and the tensile strength are listed in Table 1. As shown in Table 1, the average tensile fracture strength of the Cu/Nb joints obtained at 1253 K and 1323 K are about 184 MPa and 193 MPa, respectively. The maximum tensile strength of about 222 MPa is obtained in the Cu/Nb joint bonded at 1253 K. These tensile strengths are close to or higher than that of the pure Cu (200e300 MPa) [14]. Additionally, the average maximum tensile load and the effective fracture area of the Cu/Nb joints obtained at 1253 K are obviously larger than those of the joints obtained at 1323 K, which means that the load capability of the Cu/Nb joints obtained at 1253 K is higher than that of the joints obtained at 1323 K. Considering the results that Cu and Nb rod can't be joined at 1193 K, 1223 K, 1353 K and 1373 K, it can be concluded that the annealing temperatures between 1253 K and 1323 K are the effective temperature for the direct bonding of the Cu and Nb rods. If the annealing temperature is outside the effective temperature range 1253e1323 K, it is difficult to connect the Cu and Nb rods. Obviously, the direct bonding method of the immiscible metals is very sensitive to the annealing temperatures which need to be controlled accurately. Since the annealing temperatures are close to the melting point of Cu (TmCu) and considering that it is impossible to connect the Cu and Nb rods when the annealing temperatures are higher than TmCu (1356 K), TmCu is used as the standard to evaluate the effective temperature range for the Cu/Nb bonding system. The specific method is to calculate the ratio of annealing temperature to TmCu. So, the effective temperature range 1253e1323 K can be expressed as 0.92 TmCu-0.98 TmCu, which become a very intuitive concept. The above results can be supported by other references. For example, He and his coworkers [15] annealed an immiscible Mo/Zn bilayered sample at 673 K for 5 h under the protection of Ar gas. Since the melting temperature of Zn, TmZn, is 692.53 K, the annealing temperature of 673 K is 0.97 TmZn, which is in the effective temperature range presented in this paper. The EDX linescan profiles performed in the annealed sample show that the difficulty for the element diffusion of immiscible alloy systems is overcome and Mo atoms diffuse a long distance (~600 nm) in Zn layer at the annealing temperature.
Fig. 4. Cu/Nb joint obtained through the direct bonding method.
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Fig. 5. HRTEM analysis of the cross-section of the Cu/Nb joint bonded at 1253 K: (a) HRTEM image of the Cu/Nb interface; (b) the partially enlarged view of the region marked with white frame in (a); (c) drift-corrected spectrum image scanning; (d) the curve of diffusion depth along the red line marked in (c). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
Fig. 6. Load-Displace curves of the tensile tests of the Cu/Nb joints obtained at (a) 1253 K and (b) 1323 K, respectively.
Table 1 Tensile results of the Cu/Nb joints bonded at different temperatures. Sample Annealing temperature (K)
Maximum Effective fracture Tensile fracture tensile load (N) area (mm2) strength (MPa)
1 1253 2 1253 3 1253 Average
4507 4186 5241 4645
20.24 23.18 33.89 25.77
222.67 180.59 154.65 185.97
4 1323 5 1323 6 1323 Average
2266 3055 3025 2782
10.98 15.67 16.90 14.52
206.38 194.96 178.99 193.44
3.3. Bending strength test of Cu/Nb joints The bending strengths of the Cu/Nb joints are obtained by threepoint bending test in this study. The tests were carried out on three samples which were prepared through the same direct bonding process, where the annealing temperature was 1253 K, the bonding time was 3 h and the bonding pressure was 106 MPa. The load-displacement curves of bend tests are shown in Fig. 7 and the corresponding results are listed in Table 2. As shown in Table 2, the average bending load and bending strength of the Cu/ Nb joints can reach 836 N and 42 MPa, respectively. The above mechanical properties of the Cu/Nb joints also prove that the as-constructed Cu/Nb interfaces have high strength.
X. Pan et al. / Journal of Alloys and Compounds 723 (2017) 1053e1061
Fig. 7. Load-displacement curves of the bending tests for the Cu/Nb joints obtained at 1253 K.
Table 2 Results of the bending tests for the Cu/Nb joints obtained at 1253 K. Sample Temperature (K)
Max load (N)
Fracture area (mm2)
Max strength (MPa)
1 1253 2 1253 3 1253 Average
803.7 873.2 832.6 836.5
18.15 18.31 22.69 19.72
44.28 47.69 36.69 42.89
3.4. SEM analysis of the tensile fractures of Cu/Nb joints The SEM images shown in Fig. 8 are the tensile fractures of three Cu/Nb joint specimens annealed similarly at 1253 K for 3 h under the pressure of 106 MPa. In Fig. 8, many equiaxed dimples with different sizes and micro-holes can be obviously observed. These results indicate that the fractures of the Cu/Nb joints obtained through the direct bonding method are mainly plastic when being stretched. Fig. 8 (a)-(c) are the fractures on end-surfaces of the Cu rods of the Cu/Nb joints, where many deep dimples can be seen. Fig. 8 (d)(f) are the fractures on end-surfaces of the Nb rods of the Cu/Nb joints, where many shallow dimples can be seen. Since dimple is a peculiar feature of the fracture morphology of pure Cu, these
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results mean that the fractures occur in a part of the Cu rod near to the Cu/Nb interface, indicating that the bonding strength of the Cu/ Nb metallurgical interface is very high. In addition, at the area marked with white rectangular frame in Fig. 8 (d), a steady river pattern can be seen, which is characteristic of brittle cleavage fracture in morphology. Since the cleavage fracture usually occurs in the bcc lattice metal and Nb is just bcc lattice metal, it can be concluded that the Nb rod is fractured during the tensile at some time, which also means that the bonding strength of the Cu/Nb metallurgical interface is very high. The compositions of the dimples of the specimens have been analyzed by means of the EDX spectrometer in the SEM, which are shown in Fig. 9. Fig. 9 (a) is the fracture on the end-surface of the Nb rod of the Cu/Nb joint. Fig. 9 (b) presents the EDX line-scanning composition profile along the white line marked in Fig. 9 (a). It can be seen from Fig. 9 (b) that the two composition curves show periodic fluctuation, which are caused by the structure of the dimples. At the bottoms of the dimples, the concentration of Nb is higher than that of Cu, conversely, the concentration of Cu is higher than that of Nb at the tops of the dimples. Considering the morphology of the dimples still remain the peculiar feature of pure Cu, it can be concluded that the fracture of the Cu/Nb joints mainly occur at the end of the Cu rod. The coexistence of Nb and Cu in the dimple prove once again the mutual diffusion occurring between the Cu and Nb rods during the bonding process. 3.5. Diffusion mechanism of the direct bonding method As discussed above, the temperatures between 1253 K and 1323 K, namely 0.92 TmCu-0.98 TmCu are the effective temperature for the direct bonding of the Cu and Nb rods. Only in that effective temperature range can the mutual diffusion occurs between the Cu and Nb rods and form the metallurgical bonding interface. The reason can be described as following. In the temperature 0.92 TmCu-0.98 TmCu, Cu is in semi-solid state and the Cu may have phase transition [16,17]. In this state, 60% of Cu can be solid phase and 40% of Cu is in a molten-state [16]. Molecular dynamics simulations (MD) [18] proves that the face-centered cubic (fcc), body-centered cubic (bcc), hexagonal close-packed (hcp) structures and amorphous phase can coexist in Cu in semi-solid state [19]. And the simulation result also shows that during the processes of solidification of liquid Cu and amorphous crystallization, the crystal embryo of bcc is always formed firstly [20]. In addition, the lattice constant of Cu at 1253 K is about 0.369 nm,
Fig. 8. SEM images of tensile fracture surface corresponding to the Cu and Nb rods of three Cu/Nb joints: (a) and (d) specimen 1, (b) and (e) specimen 2 and (c) and (f) specimen 3, respectively.
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Fig. 9. Analysis of the composition of the dimples in the Nb rod of the Cu/Nb joint. (a) SEM of fracture morphology and (b) corresponding line-scanning composition profile along the white line marked in (a).
which is larger than that of Cu room temperature (0.361 nm). Correspondingly, the atomic gap of Cu will also become larger at 1253 K [21]. According to the above discussion, there may be several kinds of diffusion paths for the diffusions of Cu and Nb. Firstly, bcc structure of Cu in semi-solid state and larger atomic gap at high temperature can serve as diffusion path for the Nb atoms. Secondly, there will be some spot-defects in the Cu melt, such as vacancies [22,23]. At this time, the Nb atoms may occupy the vacancies in the Cu melt by thermal motion, reaching the purpose of diffusion. Finally, since the crystal structure of Nb is bcc, the bcc structure appears in the Cu in semi-solid state may also caused some Cu atoms diffuse into Nb. But, in general, the diffusion of Nb into Cu should be dominant because the Cu in semi-solid state can provide more diffusion paths. If the temperature is lower than 0.92 TmCu or higher than 0.98 TmCu, the diffusion path may be inadequate. The specific diffusion mechanism needs to be further explored. 3.6. Thermodynamic mechanism for the direct bonding of Cu and Nb 1. Thermodynamic model In this work, a thermodynamic model is constructed for the direct bonding method used for the connection of the immiscible Cu and Nb metals. According to the process as shown in Fig. 1, the total energy change during the bonding process can be expressed as follows:
DGtotal ¼ DGðCu=NbÞ þ DGint þ DGbulkðCu;NbÞ þ Es þ P$DV
(2)
where,
DGtotal is the total Gibbs free energy during the direct bonding process;
and DGint . On the other hand, DGðCu=NbÞ and DGint are really the
Gibbs free energy of the alloying of Cu and Nb, namely DGalloying. So, equation (2) can be rewritten as following:
DGinitial ¼ Es þ P$DV þ DGbulkðCu;NbÞ DGalloying ¼ DGðCu=NbÞ þ DGint
tween Cu and Nb, and can be described as follows: in NbÞ
Among which, DGðCu
þ DGðNb
in NbÞ
in CuÞ
(3)
is the Gibbs free energy caused by
the diffusion of Cu into Nb, and DGðNb in CuÞ is the free energy caused by the diffusion of Nb into Cu. For the solid-solution alloy system, these items are usually negative, the corresponding diffusion can occur automatically. However, for the immiscible alloy system, the items are usually positive, which need to get energy support from outside; DGint is the interfacial free energy between Cu and Nb, which
(4)
Equation (4) is the thermodynamic model constructed for the direct bonding of Cu and Nb. If DGinitial is larger than DGalloying , the diffusion and alloying between Cu an Nb can occur and proceed. In other words, the initial free energy, DGinitial can serves as the thermodynamic driving force for the diffusion and alloying at this time. The calculation formulas for the above thermodynamic quantities are based on Miedema model and Alonso method [24e26]. Generally, the values of DGðCu=NbÞ are calculated by equations (3), (5) and (6).
DGðCu
in NbÞ
¼ DH ðCu
in NbÞ
T 0 DS
(5)
DGðNb
in CuÞ
¼ DH ðNb
in CuÞ
T 0 DS
(6)
Where, DH ðCu in NbÞ /ðDH ðNb in CuÞ Þ and DS are the enthalpy and entropy of formation, respectively.
DS ¼ RðX A lnX A þ X B lnX B Þ
DGðCu=NbÞ is the Gibbs free energy caused by the diffusion be-
DGðCu=NbÞ ¼ DGðCu
also need to get energy support from outside; DGbulkðCu;NbÞ is the Gibbs free energy of the mechanical mixture of metal Cu and Nb; P$DV is the energy caused by the pressure; E s are the storage energies in the Cu and Nb rods, which could be measured by the DSC test. According to the statements above, Es , P$DV and DGbulkðCu;NbÞ form the initial Gibbs free energy of the Cu-Nb bonding system, namely DGinitial, which can provide energy support for DGðCu=NbÞ
(7)
where, R is the gas constant; XA and XB are the atomic concentrations of metal A (Cu) and B (Nb) in the alloy, respectively. According to Miedema et al., the enthalpy of formation can be calculated as follows [27e29]:
DH A
in BðB in AÞ
¼ DH cA
in BðB in AÞ
þ DH e þ DH s
(8)
where, DH c , DH e and DH s correspond to chemical, elastic and structural contributions, respectively. Since the calculated value of DH e is 6.06 104 kJ/mol in this paper and the contribution of DH s to the formation enthalpy are very small [30], the elastic and structural terms are ignored. The chemical contribution, DH c , is
X. Pan et al. / Journal of Alloys and Compounds 723 (2017) 1053e1061
1059
described as follows [31]: A in BðB in AÞ
DH cA
A in B
in B
¼ xA fAB DH sol
The solution heat of DH Cu sol Ref. [31]. 2=3
DH Asolin
B
¼
2PV A
h
gchem ¼ ss
(9) in Nb
and DH Nb sol
in Cu
are given by
2. DSC test of storage energies
(10)
The values of Es (the storage energy) in equation (4) have been measured by the DSC test according to the process described in Section 2. Before the DSC test, the Cu rod was cut into two rods with the same quality, and named as specimen 1 and specimen 2, respectively. The same process was implemented in the Nb rod, and the obtained two rods were named as specimen 3 and specimen 4. Specimen 2 and 4 were annealed at 1323 K for 6 h and 1873 K for 6 h to release the possible storage energy before the DSC tests. The final DSC results are shown in Fig. 10, where curve 1, 2 3 and 4 corresponding to specimen 1, 2, 3 and 4, respectively. It can be seen from Fig. 10 that the apparent exothermic peaks appear at curve 1 and curve 3, and the areas of the peaks are 10.36 kJ/mol and 49.30 kJ/mol. Conversely, there are no obvious exothermic peaks appearing in curve 2 and 4. Comparing the curves, it can be concluded the storage energies in the Cu and Nb rods are 10.36 kJ/ mol and 49.30 kJ/mol, respectively. It also can be seen from Fig. 10 that the storage energy in the Cu rod is released in the temperature range of 873e1073 K and the storage energy in the Nb rod is released in the temperature range of 1073e1273 K. So, when the annealing temperature is higher than 1273 K, the total released storage energy in the Cu-Nb bonding system should be the sum of the storage energies of the Cu and Nb rods.
where, P, Q and R are empirical constant; 4A and 4B are the chemical potentials of metal A and B; nwsA and nwsB are the Wigner-Seitz concentration of the cell boundary of metal A and B, respectively. The f BA is a function which accounts for the degree to which atoms of type A are surrounded by atoms of type B, given by Ref. [32].
(11)
where, s is short-range order parameter of 5 for Cu-Nb system and C SA /C SB is the surface concentration of A/B atoms. The interfacial free energy, DGint is calculated by Ref. [32].
DGint ¼ aA DGint;A þ aB DGint;B
(12)
where, aA and aB are the proportions of the interface atoms in the alloy of A and B, respectively. In our calculation, aA and aB are deduced [32].
aA ¼
PnA
DdA;i
xA Pi¼1 NA i¼1
dA;i
PnB
; aB ¼ xB Pi¼1 NB
DdB;i
i¼1
3. Thermodynamic driving forces for the diffusion and interface alloying
(13)
dB;i
Since the Cu-Nb system is immiscible, the DGbulkðCu;NbÞ term in the thermodynamic model described by equation (4) can be considered as zero. Additionally, according to Hooke's law, the volume change made by the pressure (P) in the Cu/Nb contacting surface, namely DV, is 106 m3. So, the value of the P$DV term in equation (4) is 0.32 kJ/mol. The other thermodynamic quantities including DGðCu in NbÞ , DGðNb in CuÞ and DGint are calculated ac-
The interfacial atomic free energy of 1 mol of metal is given by Ref. [32].
DGint;i ¼ S f ;i *gAB ss
(14)
s;0 s;0 þ gchem gAB ss ¼ 0:15 gA þ gB ss
(16)
2=3
C0V A
where, C 0 is a constant (4:5 108 ), and V A is the atomic volume of component A in the alloy. All the parameters used in the calculations are listed in Table 3 and Table 4.
i 1=3 1=3 2 ð4A 4B Þ2 þ QP nwsA nwsB R P 1=3 1=3 nwsA þ nwsB
2 2 C SB f BA ¼ C SA 1 þ s C SA
DH sol
(15)
cording to equation (5)e(16). Fig. 11 is the final calculated Gibbs free energy diagrams of the Cu-Nb bonding system according to the as-constructed thermodynamic model. According to Fig. 11, the green line is the alloying free energy
Where, gs;0 and gs;0 B are the surface energy of metals A and B, A being 1825 and 2700 mJ/m2 for Cu and Nb, respectively; gchem is ss given by Ref. [32].
Table 3 Calculation parameters of DGðCu=NbÞ [26]. e means the data did not exist. 2=3
metal/system
Vi
Cu Nb Cu-Nb
3.70 4.89 e
(cm2)
1=3
nws (d.u.)1/3
F (V)
Q/P
P
R (J/mol*K)
s
T0 (K)
1.47 1.68 e
4.55 4.72 e
e e 9.4
e e 14.1
e e 8.314
e e 5
e e 300
di (104cm)
Ddi (108cm)
Sf (105m2)
C0 (108)
5 11 e
e e 5
1.66 2.20 e
e e 4.5
Table 4 Calculation parameters of DGint [26]. e means the data did not exist. 2=3
metal/system
Vi
Cu Nb Cu-Nb
3.70 4.89 e
(cm2)
gs;0 (mJ*m2) i
DH Cu sol
1825 2700 e
e e 31.56
in Nb
(kJ/mol)
DH Nb sol e e 40.22
in Cu
(kJ/mol)
1060
X. Pan et al. / Journal of Alloys and Compounds 723 (2017) 1053e1061
the Cu and Nb rods completely, the diffusion and alloying in the CuNb bonding system can occur and proceed. It need to be pointed out that only a large thermodynamic driving force is not enough for the diffusion between Cu and Nb, sufficient diffusion paths are also necessary. That is the reason why the joining of the Cu and Nb rods by the direct bonding method should be carried out in a narrow and sensitive temperature range of 0.92 TmCu-0.98 TmCu. 4. Conclusions
Fig. 10. DSC curves of Cu and Nb rods: 1#-non-annealed Cu rod; 2#-annealed Cu rod; 3#-non-annealed Nb rod; 4#- annealed Nb rod.
Fig. 11. Calculated diagram of the initial energy and alloying energy of the Cu-Nb bonding system.
(DGalloying ) of the Cu-Nb bonding system and the maximum value is 19.48 kJ/mol. The red line is the initial free energy (DGinitial ) of the Cu-Nb bonding system at 873e1073 K, the maximum value of which is 10.36 kJ/mol. It can be seen that the DGinitial is much lower than the DGalloying in the most of the composition of Cu, indicating that the bonding system does not have enough thermodynamic driving force to ensure that the diffusion and alloying occur between the Cu and Nb rods. The main cause resulting in the low initial free energy is that only the storage energy in the Cu rod is released at this time. The blue line in Fig. 11 is the initial free energy (DGinitial ) of the Cu-Nb bonding system at the temperature higher than 1073 K. It can be seen that the DGinitial term is much higher than the DGalloying term, which is caused by the co-release of the storage energies of the Nb and Cu rods. The pink arrow in Fig. 11 shows the minimum difference of 40.51 kJ/mol between DGinitial and DGalloying , which indicates that the bonding system has an enough thermodynamic driving force to ensure that the diffusion and alloying occur at the Cu/Nb interface at the temperature higher than 1073 K. Since the calculated P$DV is very small and the DGbulkðCu;NbÞ term is zero, the initial energy, DGinitial depends almost entirely on the release of the storage energies (Es ) in Cu and Nb rods. In other words, Es serves as the thermodynamic driving force for the diffusion and alloying occurring at the Cu/Nb interface during the direct bonding process. Only the Cu-Nb bonding system is annealed at sufficiently high temperature to release the storage energies in
In this paper, a direct bonding method was used to construct the metallurgical interface directly between immiscible Cu and Nb without using an interlayer metal. The construction processes include the polish for the end-surfaces of the Cu and Nb rods, the coaxial assembly of the Cu and Nb rods, and sequential annealing at high temperature for 3 h under a pressure of 106 MPa. Through the method, a Cu/Nb metallurgical interface was successfully constructed and the corresponding Cu/Nb joint was obtained. The maximum tensile strength and bending strength of the obtained Cu/Nb joints are about 222 MPa and 47 MPa, respectively, which have reached a very high level. The micro-test results show that thickness of the Cu/Nb diffusion layer is about 36 nm. Additionally, a thermodynamic model was established for the direct bonding of the immiscible metals by using the Miedema theory. Through the calculations based on the model and differential scanning calorimetry (DSC) tests, the storage energies of Cu and Nb rods released at high temperature are confirmed to be able to overcome the positive reaction heat of Cu-Nb system, and serve as the thermodynamic driving force for the diffusion and alloying occurring at the Cu/Nb interface. Finally, it need to be pointed out that the key point of the direct bonding method is the annealing temperature, which should be close to the melting point of Cu (TmCu ¼ 1356 K). On the one hand, the high-temperature structure of Cu at the temperature close to the TmCu can provide the diffusion paths for Cu and Nb probably. On the other hand, the storage energies of Cu and Nb rods can be fully released to drive the diffusion. The experimental results show that the suitable annealing temperature should be between 0.92 TmCu0.98 TmCu, which is very sensitive and need to be controlled accurately. Acknowledgement This work was supported by the National Natural Science Foundation of China (51171118 and 51471114), and the Science and Technology Support Project of Tianjin City (11ZCKFGX03800). References [1] X. Zhang, N. Li, O. Anderoglu, H. Wang, Nuclear instruments and methods in physics research section B: beam interactions with materials and atoms, Nucl. Instrum. Methods Phys. B 261 (2007) 1119. [2] A. Misra, M.J. Denkowicz, The radiation damage tolerance of ultra-high strength nanolayered composites, JOM 59 (2007) 62e65. [3] W.Z. Han, A. Misra, N.A. Mara, Role of interfaces in shock-induced plasticity in Cu/Nb nanolaminates, Philos. Mag. 91 (2011) 4172. [4] PonoB.И.A., ШИPЯCBa H.B. // ЖypHaл HeopraHHqeckoй Xиmии. T.6. 10. (1961) 2334e2340. [5] W. Elthallabawy, T. Khan, Liquid phase bonding of 316L stainless steel to AZ31 magnesium alloy, Mater. Sci. Technol. 27 (2011) 22e28. [6] M. Ding, S.S. Liu, Y.C. Wang, TIG-MIG hybrid welding of ferritic stainless steels and magnesium alloys with Cu interlayer of different thickness, Mater. Des. 88 (2015) 375e383. [7] F.M. Zhou, F.Q. Zhang, F.Y. Song, Vacuum diffusion bonding and interfacial structure in Ta/Cu couple, Rare Metal Mater. Eng. 42 (2013) 1785e1789. [8] S. Wang, Y.H. Ling, J.J. Wang, Microstructure and mechanical properties of W/ Cu vacuum diffusion bonding joints using amorphous Fe-W alloy as interlayer, Vacuum 114 (2015) 58e65.
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