Shen-Hung Wei and Chien-Cheng Lina)
Department of Materials Science and Engineering, National Chiao Tung University, Hsinchu 300, Taiwan (Received 18 November 2013; accepted 24 January 2014)
In this study, 3 mol% Y2O3-stabilized zirconia (3Y–ZrO2) and commercially pure titanium (cp-Ti) joints were fabricated with an Ag68.8Cu26.7Ti4.5interlayer (Ticusil) at 900 °C for various brazing periods. After brazing at 900 °C/0.1 h, Ti2Cu, TiCu, Ti3Cu4, and TiCu4layers were present at the Ti/Ticusil interface, while TiCu and TiO layers were observed at the Ticusil/3Y–ZrO2interface. In the residual interlayer, clumpy TiCu4was formed along with the Ag solid phase. After brazing at 900 °C/1 h, Ti3Cu3O and Ti2O layers were formed at the interlayer/ZrO2interface, while Cu2O was precipitated in the residual interlayer with 111½ Cu2O== 111½ Agand 20ð 2ÞCu2O== 202ð ÞAg. After brazing at 900 °C/6 h, a two-phase (a-Ti 1 Ti2Cu) region was observed on the Ti side with
2110
½ aTi== 100½ Ti2Cuand 0002ð ÞaTi== 013ð ÞTi2Cu, while the TiCu layer grew at the expense of
Ti3Cu4and TiCu4. The bonding mechanisms and diffusion paths were explored with the aid of Ag–Cu–Ti and Ti–Cu–O ternary phase diagrams.
I. INTRODUCTION
The Ag–Cu eutectic alloy with 28Cu (wt%) is frequently used in the conventional brazing of Ti alloys and other materials at temperatures higher than 800 °C.1The brazed assembly exhibits good bonding strength, resulting from the formation of Ti–Cu compounds that enhance the wettability of the solid/liquid interface.2,3 However, Sn, Ag, Cu, Ni, Co, and Fe exhibit very poor wettability on the surface of ceramics.4 The wettability can be significantly improved by adding active metals, such as Ti, Zr, and Hf, to the brazing filler. These metals can effectively reduce the interfacial energy by their strong chemical attraction to the oxide ceramic and can enhance the wetting behavior,5 which even enables the surface of the ceramic to be wetted by the formation of the intermediate layer during brazing and provides a good joint.6,7
Hanson et al.7 observed that a small percentage of Ti addition (;4.5Ti wt%) to the braze alloy led to a significant increase in the four-point bend strength of a PSZ–PSZ (PSZ, partially stabilized zirconia) joint. In comparison with PSZ–PSZ or PSZ-stainless steel joints, the 3Y–ZrO2/Ti brazing couple using a Cusil ABA [63Ag–35.25Cu–1.75Ti (wt%)] had a superior joint strength due to the suitable reaction layers at the interface. The microstructural evolu-tion at the interface between the Ti alloy (or ZrO2) and braze alloy has been previously explored based on microstructural characterization using scanning electron microscopy (SEM)/energy dispersive spectroscopy (EDS).1–3,6,8 However, the SEM/EDS spatial resolution was limited to
the micro range (,2 lm) and unable to analyze the crystal structures. In this study, the brazefiller of Ticusil [68.8Ag–26.7Cu–4.5Ti (wt%)] was selected to join the 3Y–ZrO2and Ti at 900 °C for various brazing periods. The objective was to elucidate the microstructural evolution and bonding mechanisms of the sandwiched Ti/Ticusil/3Y–ZrO2 joints based on the microstructural characterization using SEM/EDS and transmission elec-tron microscopy (TEM)/EDS.
II. EXPERIMENTAL
Hot pressed 3 mol% Y2O3-stabilized ZrO2 bulks [3Y–ZrO2, 3 mm thickness on average, with an apparent density of 5.9 g/cm3, mainly .94ZrO2 1 HfO2, ,5.4Y2O3(wt%), Toyo Soda Mfg. Co., Tokyo, Japan] and commercially pure titanium billets [cp-Ti, 5 mm thickness on average, mainly 99.31Ti–0.30Fe–0.25O (wt%), Kobe Steel, Ltd., Tokyo, Japan] were used as starting materials. The 3Y–ZrO2 and Ti were brazed by a Ticusil interlayer [50lm thick, 68.8Ag–26.7Cu–4.5Ti (wt%), Wesgo, Inc., Belmont, CA] under 0.1 MPa at 900 °C for 0.1, 1, and 6 h in a tube furnace (Model STF54434C, Lindberg/Blue, Thermo Elec. Inc., TX). The process temperature (900 °C) was selected to be 70 °C above the eutectic point of thefiller. To prevent the sample from being oxidized, the working chamber was preevac-uated to 1 104Pa and purged with Ar (99.99% pure) to 1 atm for at least four cycles. The heating and cooling rates were 5 and 3° C/min, respectively.
Cross-sectional SEM specimens were cut, ground, and finally polished using conventional standard procedures. Then the microstructures of various reaction zones were
a)
Address all correspondence to this author. e-mail: [email protected] DOI: 10.1557/jmr.2014.30
characterized using a SEM (Model JSM 6500F, JEOL Ltd., Tokyo, Japan) operating at 15 kV, which was equipped with an EDS (Model ISIS 300, Oxford Instrument Inc., London, U.K.). Cross-sectional TEM specimens were prepared using a focused ion beam (FIB, Model Nova 200, FEI Co., Hillsboro, OR) and the lift-out technique.9 Typical TEM specimens had dimensions of approximately 17 lm in width and 2 lm in depth. FIB milling was performed with a Ga1ion beam at 30 keV. Before the FIB milling, a Pt layer approximately 1 lm thick was deposited by ion beam chemical deposition using C9H16Pt as the precursor gas. The Pt layer served two main purposes: being a marker illustrating the location and protecting the sample from being directly exposed to Ga1ion beam implantation. After rough milling (7–1 nA), polishing (0.5–0.1 nA), and final polishing (50–10 pA), a thin foil (17 2 0.05 lm) was cut off, and then, a micromanipulator was used to transfer the foil from the sample to a TEM grid (Formvar/carbon coated-copper 200 mesh) for subsequent TEM analyses.
The interfacial microstructures of the jointed samples were then observed using an analytical TEM (Model JEM 2000 FX, JEOL Ltd, Tokyo, Japan). The Cliff– Lorimer standardless technique was performed to analyze the chemical compositions of various phases using an ultra-thin window energy dispersive spectrometer (EDS, ISIS 300, Oxford Instrument Inc., London, U.K.) attached to the TEM.10,11Based on the thin-film criterion, assuming that the effects of absorption and fluorescence could be neglected, all the measurements were conducted on the Ka lines of corresponding elements from the very thin regions of the TEM specimens. The standardless k factors in the Cliff–Lorimer equation had the same precision as the ones measured by the standard approach if the thin-film criterion of Ka lines was fulfilled.11 The crystal structures and lattice parameters were determined from analysis of the selected area diffraction patterns (SADPs) using the computer-simulation software of crystallography (CaRIne Crystallography 3.1, DIVERGENT S. A., Compiègne, France).
III. RESULTS AND DISCUSSION A. SEM/EDS analyses
Figure 1 presents a backscattering electron image (BEI) of the as-received Ticusil foil after annealing at 900 °C for 1 h. The clumpy Ag-rich phase was observed in addition to the lamellar eutectic structure of (Ag) 1 (Cu) within the Ticusil foil. In addition, the Ti3Cu4phase [42.5Ti–56.7Cu–0.8Ag (at.%)] was also observed. It was believed that the reactions between Ti and Cu pushed the Ag–Cu composition into the hypoeutectic region, result-ing in the formation of a clumpy proeutectic Ag-rich phase and the lamellar eutectic. The eutectic lamellae
consisted of the bright Ag-rich phase and the dark Cu-rich phase. Furthermore, because the solubility of Cu in Ag was limited,12small spherical Cu-rich particles were precipitated from the Ag phase during subsequent cooling, which was consistent with the observations by Lin et al.13
Figure 2 presents the BEIs of the interfacial microstruc-ture in between Ti and Ticusil (or Ticusil and 3Y–ZrO2)
FIG. 1. SEM BEI microstructure of the as-received Ticusil foil before
joining and after annealing at 900 °C for 1 h, showing the Ag–Cu
eutectic, Ti3Cu4phase, and Ag solid phase with Cu inclusions.
FIG. 2. SEM BEI microstructure of the Ti/Ticusil/3Y–ZrO2joint after
brazing at 900 °C for (a) 0.1 h, (b) 1 h, and (c) 6 h, showing the reaction
layers at the Ti/Ticusil interface [Ti2Cu (A), TiCu (B), Ti3Cu4 (C),
TiCu4(D), anda-Ti 1 Ti2Cu (I)], the reaction layers at the Ticusil/
3Y–ZrO2interface [TiCu (E), TiO (F), Ti3Cu3O (G), and Ti2O (H)],
and the residual Ticusil interlayer [TiCu4(D), Cu2O, Ti3Cu3O, and
after brazing at 900 °C for various brazing periods. The original Ti/Ticusil and Ticusil/3Y–ZrO2 interfaces were marked by arrows. The dissolution of the Ti substrate into the liquid interlayer was thought to be much more prominent compared with that of the 3Y–ZrO2 substrate. Figure 2(a) reveals the BEI of the Ti/Ticusil/3Y–ZrO2joint after brazing at 900 °C for 0.1 h. Four distinct interfacial layers (marked A, B, C, and D) existed between Ti and Ticusil. The relative contrast of the interfacial layers reflected a decreasing Ti/Cu ratio from the left to right sides because of the atomic number effect on the BEI. The thinner layer A was tentatively identified as a Ti2Cu layer. Layer B consisted of 47.7Ti– 48.6Cu–3.7Ag (at.%), indicating that layer B was TiCu. Layer C contained 41.4Ti–57.0Cu–1.6Ag (at.%) adjacent to the original interface between Ti and Ticusil, corresponding to the Ti3Cu4phase. A dendritic TiCu4phase (marked D), consisting of 21.3Ti–77.2Cu–1.5Ag (at.%), tended to grow into the residual interlayer. A clumpy phase (marked D, TiCu4) was observed together with the Ag solid phase within the residual interlayer. In addition, the reduction of 3Y–ZrO2 by the Ticusil interlayer7,14led to the formation of a dark layer (F), consisting of 49.3Ti–48.9O–1.8Cu (at.%), which corresponds to the TiO phase. The dissolved Ti reacted with Cu to form a relatively gray layer (marked E), corresponding to the TiCu phase, near the TiO layer in the residual interlayer. The TiO layer existed as a diffusion barrier layer that limited the migration of Zr into the residual interlayer.15 The consumption of Cu atoms due to the formation of various Ti–Cu compounds in the molten interlayer resulted in the deviation of the residual interlayer from the Ag–Cu eutectic toward the Ag-rich region.16Therefore, the Ag solid phase was observed throughout the residual interlayer after cooling. The Ag solid phase in the residual interlayer consisted of 94.4Ag–2.0Ti–3.6Cu (at.%) based on the EDS analyses. In this case, the yttrium (the stabilizer of ZrO2) was not detected in TiO, indicating that the original interface was located between TiO and 3Y–ZrO2. Figure 2(b) displays the reaction layer sequence of a-Ti/Ti2Cu(A)/TiCu(B)/Ti3Cu4(C)/TiCu4(D)/residual interlayer/Ti3Cu3O(G)/Ti2O(H)/TiO(F)/ZrO2x in the Ti/Ticusil/3Y–ZrO2joint after brazing at 900 °C for 1 h. The fact that layer Ti2Cu3was missing between Ti3Cu4 and TiCu4 could be explained by kinetics considera-tions.17 The residual interlayer contained the Ag solid phase, withfine Cu2O precipitates replacing the clumpy TiCu4phase. Two oxide layers (G and H) were observed to abut the TiO layer, indicating that an extensive reduction of ZrO2to oxygen-deficient zirconia occurred due to the dissolved Ti in the molten interlayer. Further-more, O exhibited a higher reactivity with Cu than Ag such that the Cu2O phase was formed in the Ag solid phase during cooling. According to the Ti–Cu–O ternary phase diagram,18 the Ti–Cu–O compound would be formed while exceeding the maximum solubility of O (;2 at.%) in the TiCu. For comparison, the reaction zone
between the Ag–Cu–Sn–Ti filler metal and Al2O3 was comprised of two sublayers after brazing at 900 °C for 20 min: an fccc-TiO layer and another sublayer of Ti3Cu3O with a diamond cubic structure, which provided a more reliable joint than TiO alone.19Figure 2(c) illustrates the reaction layer sequence ofa-Ti/a-Ti 1 Ti2Cu(I)/Ti2Cu(A)/ TiCu(B)/residual interlayer/Ti3Cu3O(G)/Ti2O(H)/TiO(F)/ ZrO2x in the sandwiched joint after brazing at 900 °C for 6 h, where the Cu2O, Ti3Cu3O, and Ag phases were present in the residual interlayer. Cu diffusion from the interlayer to the Ti substrate led to the hypoeutectoid transformation of a two-phase layer (a-Ti 1 Ti2Cu), designated as I, during a long-term brazing period.20The elongated a phase in the (a-Ti 1 Ti2Cu) region corre-sponded to a proeutectoid phase. The transformation could be written by two steps of the hypoeutectoid reaction: b-Ti ! b1-Ti 1 a-Ti (proeutectoid a-Ti) and b1-Ti ! Ti2Cu 1 a-Ti (eutectoid a-Ti). The formation Gibbs free energies (ΔG0) of Ti2Cu and TiCu were more negative than those of Ti3Cu4and TiCu4.21Therefore, it was believed that a further reaction could occur between Ti3Cu4(or TiCu4) and Ti to produce the more stable TiCu phase.
B. TEM/EDS analyses
1. Microstructure at the Ticusil/3Y–ZrO2interface Figure 3(a) presents a brightfield image (BFI) of two suboxide layers, corresponding to TiO and Ti2O (labeled F and H) in the Ti/Ticusil/3Y–ZrO2joint after brazing at 900 °C/1 h. Figures 3(b) and 3(c) display the SADPs of the b-TiO with the incident electron beam along the zone axes of 001½ and 111½ , respectively, indicating that the material had a cubic structure similar to the NaCl structure, with the measured lattice parameter a 5 0.4234 nm (b-TiO, Fm3m, a 5 0.4185 nm in JCPDS No. 772170).22Note that the reflections should be absent due to the zero structure factor for such planes as (010), (100), (101), and (110) but appeared in the pattern due to double diffraction and the nonstoichiometric composition. Figures 3(d) and 3(e) present the SADPs of Ti2O on the zone axes of ½1210 and 2423½ , respectively, identified as a hexagonal structure with the measured lattice parameters a 5 0.2898 nm and c 5 0.48765 nm (Ti2O, P3m1, a5 0.2953 nm, c 5 0.48454 nm in JCPDS No. 731582).23 Figure 3(f) demonstrates that the b-TiO consisted of 49.65Ti–47.33O–1.76Zr–1.26Cu (at.%). Figure 3(g) pres-ents the EDS results of Ti2O, indicating that the material consisted of 65.46Ti–30.3O–3.12Zr–1.12Cu (at.%).
ZrO2 tends to be reduced to an oxygen-deficient zirconia (ZrO2x) in a reducing environment.24 Because Ti has a high affinity for O, it can act as a reducing agent of ZrO2. When Ti is in contact with ZrO2, oxygen can diffuse out of ZrO2 and be dissolved in Ti, leading to the formation of ZrO2x, even though the Ellingham diagram
indicates that ZrO2is more thermodynamically stable than TiO2. Thus, ZrO2 is not completely reduced into the metallic Zr, while Ti is not oxidized as TiO2in the partial reduction and dissolution process mentioned above. However, certain titanium oxides are likely formed as the amount of dissolved oxygen increases in Ti. In this study, ZrO2 was partially reduced by Ti to become ZrO2x, while oxygen was dissolved into the residual interlayer to form the intermediate compounds TiO, Ti2O, and Ti3Cu3O. The liquid Ticusil exhibited good wettabil-ity on the 3Y–ZrO2surface due to the formation of TiO (metallic characteristics).25
2. Microstructure at the residual Ticusil interlayer Figure 4(a) presents a BFI of the Cu2O precipitate in the Ag-rich matrix within the residual interlayer after brazing at 900 °C for 1 h. The Ag-rich matrix was accompanied by dislocations because of the CTE (coefficient of thermal expansion) mismatch between Ticusil and 3Y–ZrO2. Figure 4(b) displays the SADPs of the Cu2O and Ag-rich phase with its schematic diagram redrawn in Fig. 4(c). From their superimposed SADPs, the orientation rela-tions of Cu2O and Ag were identified as follows:
111
½ Cu2O== 111½ Agand 20ð 2ÞCu2O== 202ð ÞAg. The lattice
parameter of Cu2O was measured to be 0.4264 nm (Cu2O, Pn3m, a5 0.4267 nm in JCPDS No. 782076). Furthermore, the lattice parameter of fcc Ag was measured to be 0.4198 nm (Ag, Fm3m, a5 0.4086 nm in JCPDS No. 040783). The orientation relationship between these two phases was cube-on-cube. Because Cu2O is an antifluorite derivative with 4 Cu atoms at 0, 0, 0 1 face-centering
FIG. 3. TEM results of the microstructure: (a) morphology ofb-TiO and Ti2O at the Ticusil/3Y–ZrO2interface of the Ti/Ticusil/3Y–ZrO2joint after
brazing at 900 °C for 1 h; (b and c) electron diffraction patterns ofb-TiO with the 001½ and 111½ zone axes, respectively; (d and e) electron diffraction
patterns of Ti2O with the ½1210 and 2423½ zone axes, respectively; (f ) EDS spectrum of b-TiO; and (g) EDS spectrum of Ti2O.
FIG. 4. TEM results of the microstructure: (a) morphology of the Cu2O
precipitates in the Ag-rich matrix of the residual Ticusil interlayer after brazing at 900 °C for 1 h and (b and c) electron diffraction patterns of
Cu2O and Ag with their schematic diagram (●, Cu2O and , Ag)
translations and 2 O atoms at 1/4, 1/4, 1/41 body-centering translations,26the (hkl) reflection spots were absent for two even and one odd integer.
3. Microstructure at the cp-Ti/Ticusil interlayer Figures 5(a) and 5(b) display BFIs of the Ti/Ticusil interfacial microstructure after brazing at 900 °C/1 h, revealing several phases, such as a-Ti, Ti2Cu, TiCu, TiCu4, and the Ag-rich matrix. Figure 5(c) shows the SADP of Ti2Cu along the zone axis of 010½ , corre-sponding to a tetragonal structure with the measured lattice parameters a 5 0.2941 nm and c 5 1.0779 nm (Ti2Cu, I4/mmm, a5 0.29438 nm, c 5 1.07861 nm in JCPDS No. 720441). Figure 5(d) presents the SADP of TiCu along the zone axis of 0½ 21, identified as a tetrag-onal structure with the measured lattice parameters a 5 0.323 nm and c 5 0.6132 nm (TiCu, P4/mmm, a5 0.3108 nm, c 5 0.5887 nm in JCPDS No. 070114). Figure 5(e) presents the EDS results of Ti2Cu, consisting of 67.62Ti–31.25Cu–1.13Ag (at.%). Figure 5(f) presents the EDS results of TiCu, which consisted of
48.13Ti–47.59Cu–4.28Ag (at.%). Figure 5(g) shows that the TiCu4contained 20.05Ti–78.60Cu–1.35Ag (at.%).
Figure 6(a) demonstrates that acicular Ti2Cu was precipitated on the titanium side of the Ti/Ticusil/3Y– ZrO2joint after brazing at 900 °C/6 h. Figure 6(b) displays the SADPs, with their patterns redrawn in Fig. 6(c), ofa-Ti and Ti2Cu with the following orientation relationship:
2110
½ aTi== 100½ Ti2Cu and 0002ð ÞaTi== 013ð ÞTi2Cu. The
Ti2Cu had a tetragonal structure with the measured lattice parameters a5 0.297 nm and c 5 1.071 nm (Ti2Cu, I4/mmm, a5 0.29438 nm, c 5 1.07861 nm in JCPDS No. 720441), while the matrix was identified as a-Ti with the measured lattice parameters a5 0.299 nm and c 5 0.474 nm (a-Ti, P63/mmc, a5 0.295 nm, c 5 0.4682 nm in JCPDS No. 441294).
C. Microstructure development and diffusion path
1. Diffusion path for the Ag–Cu–Ti phase diagram Figure 7(a) presents an isothermal section of the Ag– Cu–Ti ternary phase diagram at 900 °C.27The diffusion path (solid line), indicative of the compositions along the
FIG. 5. TEM results of the microstructure: (a and b) morphologies of thea-Ti, Ti2Cu, TiCu, TiCu4, and Ag-rich phases at the Ti/Ticusil interface of
the Ti/Ticusil/3Y–ZrO2joint after brazing at 900 °C for 1 h; (c) electron diffraction patterns of Ti2Cu with the 010½ zone axis; (d) electron diffraction
longitudinal direction perpendicular to the interface, was proposed to be a–b–c–d–e–f–g–h–i–j for the Ti/Ticusil interface after brazing at 900 °C/0.1 h. The diffusion path crossed thefields of a-Ti, a-Ti 1 Ti2Cu, Ti2Cu, Ti2Cu1 TiCu, TiCu, TiCu1 Ti3Cu4, Ti3Cu4, Ti3Cu41 L, and L. The diffusion path did not pass through Ti2Cu3from the viewpoint of kinetics.17 The three-phase regions and crossing parallel tie lines (not shown) in the two-phase regions of the ternary phase diagram were simply in correspondence to the layer interface. Figure 7(b) presents the microstructures that developed in the Ti/Ticusil joint upon brazing at 900 °C/0.1 h. The connecting letters indicate the relationship between the microstructure of the Ti/Ticusil interface and the isothermal section of the Ag–Cu–Ti phase diagram. As illustrated in Fig. 7(b), the regions ofa-Ti 1 Ti2Cu (b to c), Ti2Cu1 TiCu (d to e), TiCu1 Ti3Cu4( f to g), and Ti3Cu41 L (h to i) correspond to the interfaces betweena-Ti and Ti2Cu, Ti2Cu and TiCu, TiCu and Ti3Cu4, and Ti3Cu4and L, respectively. Thus, the layers ofa-Ti, Ti2Cu, TiCu, Ti3Cu4, and L (Ag (major)1 TiCu4 (minor)) were formed in sequence from Ti to Ticusil upon cooling.
2. Diffusion path for the Ag–Cu–Ti phase diagram Figure 8(a) illustrates the correlation between the micro-structure of the Ticusil/3Y–ZrO2interface and the Ti–Cu–O ternary phase diagram.18The diffusion path lying upon the solid line was designated as follows: a9–b9–c9–d9–e9–f9. The crossed fields were TiO, TiO 1 Ti2O, Ti2O, Ti2O 1 Ti3Cu3O, and Ti3Cu3O. The interfaces of various reaction layers were illustrated by these regions, such as TiO1 Ti2O (b9 to c9) and Ti2O1 Ti3Cu3O (d9 to e9). The reaction layers of TiO, Ti2O, and Ti3Cu3O, as illustrated in Fig. 8(b), were formed between Ticusil and 3Y–ZrO2 after brazing at 900 °C/1 h. No Ti3O2phase appeared because the peritectoid reaction (a-Ti 1 a-TiO ! Ti3O2) occurred at 920 °C.
D. Proposed model of microstructural evolution
1. Melting and solidification of Ticusil
Figure 9 shows the liquidus surface projection of the ternary Ag–Cu–Ti diagram (middle)28 and the relevant
binary phase diagrams (Ti–Cu, Ag–Cu, and Ag–Ti).12,29,30 While there are two intermetallics (Ti2Ag and TiAg) in the Ag–Ti phase diagram,29 six intermetallics (Ti2Cu, TiCu, Ti3Cu4, Ti2Cu3, TiCu2, and TiCu4) exist in the Ti–Cu phase diagram30due to the high negative values for the heat of mixing. Thus, Ti has a greater affinity for Cu than for Ag. The Ticusil alloy has a composition of 55.4Ag–36.5Cu– 8.1Ti (at.%), which might be considered as a eutectic Ag60Cu40 modified with 8.1Ti (at.%). Its composition (labeled M) fell inside a wide miscibility loop in the ternary Ag–Cu–Ti liquidus projection. Two arrows
FIG. 6. TEM results of the microstructure: (a) morphology of acicular Ti2Cu precipitates in thea-Ti matrix on the Ti side of the Ti/Ticusil/3Y–ZrO2
joint after brazing at 900 °C for 6 h and (b and c) electron diffraction patterns ofa-Ti and Ti2Cu with their schematic diagram (●,a-Ti and , Ti2Cu)
showing z¼ 2110½ aTi== 100½ Ti2Cuand 0002ð ÞaTi== 013ð ÞTi2Cu.
FIG. 7. (a) An isothermal section of the Ag–Cu–Ti system at 900 °C
and the schematic diffusion path: a–b–c–d–e–f–g–h–i–j and (b) the
correlation between the ternary phase diagram and the microstructural
evolution at the Ti/Ticusil interface of the Ti/Ticusil/3Y–ZrO2joint after
brazing at 900 °C for 0.1 h. L1and L2are represented as an Ag–Cu rich
pointed to the compositions of Laand Lbupon reaching the miscibility boundary line, indicating that the Ticusil was divided into two immiscible liquids: M! La1 Lbat 900 °C. Estimated from the miscibility loop, the liquid Ticusil was separated into 20% La[7.1Ag–59.3Cu–33.6Ti (at.%)] and 80% Lb[67.7Ag–31Cu–1.3Ti (at.%)]. Lawas approximate to the apparent ratio Ti:Cu5 1:2, while the chemistry of Lbwas close to the Ag–Cu eutectic compo-sition [60Ag–40Cu (at.%)] at 780 °C.
Although the concentrations of Ti were quite different in both the Cu–Ti rich (La) and Ag–Cu rich (Lb) liquids at 33.6 and 1.3 (at.%), respectively, the two separated liquids had the same activities of Ti and Cu, respectively. Thus, Ag could enhance the activity of Ti, while Cu suppressed the activity of Ti in the liquid.31 Because Ticusil was regarded as a modified eutectic 60Ag–40Cu (at.%) with a low content of Ti, a vertical Cu–Ti section of the Ag–Cu–Ti system [fixed 60Ag (at.%)] would be useful to analyze the solidification of the interlayer.27
The compositions of these two separated liquid phases were fixed with a small variation in their relative amounts as the point M was projected to the point X, in which the composition X was located at 60Ag–35Cu–5Ti (at.%). Figure 10(a) indicates that the cooling path of the point X would pass through several domains upon cooling from 900 °C as follows: (i) La1 Lb(two liquid phases separated at 900 °C); (ii) La1 Lb 1 Ti3Cu4 (La exhausted and transformed into Ti3Cu4 with all the Ag (La) and excess Cu (La) diffusing from
La to Lb at 889–879 °C); (iii) L 1 Ti3Cu4 (Lb trans-formed into L at 879–854 °C); (iv) L 1 (Ag) 1 Ti3Cu4 (a solid Ag phase precipitated from L at 854–843 °C); (v) L1 Ti2Cu31 (Ag) (all the Ti3Cu4transformed into Ti2Cu3and more Ag precipitated at 843–808 °C); (vi) L 1 (Ag) 1 TiCu4 (all the Ti2Cu3 transformed into TiCu4, and more Ag precipitated at 808–783 °C); and (vii) (Ag)1 (Cu) 1 TiCu4(L consumed completely and transformed into Ag and Cu below 783 °C). If this description was true, the final microstructures of the Ticusil interlayer should have consisted of Ag and Cu solid solutions and TiCu4 precipitates. However, primary Ag with Cu inclusions, eutectic (Ag)1 (Cu), and Ti3Cu4 were observed in the as-received Ticusil interlayer. The vertical Cu–Ti sections of the Ag–Cu–Ti system forfixed 55.4 and 60 at.% Ag were significantly different such that the relevant cooling path should be somewhere between M and X. Therefore, the cooling path couldfirst enter the (L 1 TiCu 1 Ti3Cu4) domain, resulting in the isothermal reaction at 900 °C in the Ag–Cu–Ti triangle, as illustrated in Fig. 9: L}
1þ TiCu4L}2þ Ti3Cu4 (L}1 and L}2 were close to La and Lb, respectively).28Thus, not only two separated liquids (La 1 Lb) but also two Ti–Cu phases (TiCu 1 Ti3Cu4) existed at 900 °C. In brief, the amount of La was consumed with the decreasing temperature, and the remaining Lb (designated L hereafter) would follow the cooling paths from U4 toward U5, as shown in Fig. 10(b), in the temperature range from 900 to 843 °C as follows17: U4 : L þ TiCu ! Ag þ Ti3Cu4 and U5: L þ Ti3Cu4 ! Ag þ Ti2Cu3. The solidus point of Ticusil (830 °C) corresponded to the transformation of Ti3Cu4to Ti2Cu3. The solidification of the liquid phase would be completed at U5. The Ti3Cu4 phase was retained at room temperature after cooling from 900 °C because the formation of Ti2Cu3had very slow kinetics.17 Thus, most of the Ti in the as-received Ticusil interlayer was associated with the Cu-rich phase (Ti3Cu4) at 900 °C.
2. Microstructural evolutions at the Ti/Ticusil/ 3Y–ZrO2interface
The model of melting and solidification mentioned above could not be applied to brazing because the reactive wetting and interdiffusion between the Ticusil interlayer melt and base materials caused the formation of Ti–Cu compounds and a dramatic compositional variation in the liquid interlayer. Based on the activity distributions of Ti and Cu, the microstructural development and bonding mechanisms of the Ti/Ticusil/3Y–ZrO2 joint were described as follows.
In stage 1, the Ticusil interlayer was melted at 900 °C, resulting in the phase separation of La [7.1Ag–59.3Cu– 33.6Ti (at.%)] and Lb[67.7Ag–31Cu–1.3Ti (at.%)]. These
FIG. 8. (a) An isothermal section of the Ti–Cu–O system at 945 °C and
the schematic diffusion path: a9–b9–c9–d9–e9–f9 and (b) the correlation
between the diffusion path and the ternary phase diagram showing the
microstructural evolution at the Ticusil/3Y–ZrO2interface of the Ti/
two phases were in equilibrium such that the activities of the Ti and Cu in them werefixed, respectively, through the interlayer, as shown in Fig. 11(a). Because the Ti–Cu affinity was stronger than the Ag–Ti affinity, the near Ti: Cu5 1:2 stoichiometry liquid (La) tended to approach the Ti side, while the near Ag–Cu eutectic liquid (Lb) moved to the 3Y–ZrO2 side. In stage 2, the Ticusil melt had a good wettability on the Ti surface because Ti was a very active metal. Good wettability could enhance the dissolu-tion and reacdissolu-tion of Ti and 3Y–ZrO2 into the Ticusil liquid. After the Ti substrate was significantly dissolved in the Ticusil interlayer, the Ti activity increased, and a gradient was established between the Ti and 3Y–ZrO2. When the Ti activity was sufficiently high in the liquid
(La) adjacent to the Ti surface, Ti3Cu4began to nucleate and grow as a thin outer layer on the Ti surface. Once Ti3Cu4was formed, the concentration of Cu adjacent to the Ti substrate was reduced, leading to a Cu activity gradient in the liquid phase abutting the Ti, as shown in Fig. 11(b). In stage 3, the predominant mechanisms were interdiffusion and a chemical reaction between Ti and the Ticusil melt. It was believed that Cu diffused across the Ti3Cu4 layer and reacted with Ti to produce TiCu and Ti2Cu at the Ti/Ti3Cu4 interface because the Cu atoms diffused faster than the Ti atoms. Figure 11(c) shows that the activities of Ti and Cu descended in opposite directions through the three Ti–Cu layers. As the Ti–Cu layers increased in thickness, the increasing distance over which
FIG. 9. Liquidus surface projection onto the Ag–Cu–Ti compositional triangle (middle), Ti–Cu binary phase diagram (top), Ag–Cu phase
diagram (left), and Ag–Ti phase diagram (right). Designations: ― (boundary of miscibility loop); !. . . [vertical Cu–Ti–(60% Ag)]; ★!
(composition of Ticusil and liquid phase separation path); M, nominal composition of Ticusil (Ag55.4–Cu36.5–Ti8.1); and, composition of
the Cu and Ti needed to diffuse hindered the transportation of Ti through the Ti–Cu layers. Thus, the further supply of Ti atoms to the interface was frequently negligible and insufficient, resulting in a Ti-depleted zone in liquid La abutting the Ti substrate after the nucleation of TiCu4 occurred at the solid/liquid interface with a reduced relative Ti/Cu ratio (3/4 ! 1/4). This result indicated that Ti was sharply consumed and that concentration inversion occurred. In stage 4, all the Ti in the liquid La was nearly consumed. Meanwhile, the TiCu4tended to grow into the liquid with more concentrated Ti and Cu (Lb) such that dendritic TiCu4 would nucleate and grow into the liquid phase. The clumpy TiCu4was also formed far away from the original Ti/Ticusil interface before the complete solidification of the interlayer. The residual interlayer hereafter could be considered an Ag–Cu binary alloy (L) with a small amount of Ti in solution. However, the Cu content in the L was gradu-ally reduced due to the formation of Cu–Ti intermetallic compounds as the composition of the liquid phase varied along the liquidus. Thus, the melting temperature would be increased somewhere in the Ag–Cu binary phase diagram.12It was believed that the composition of the residual Ag–Cu alloy (L) should be below the hypoeutectic limit (14 at.%) because no eutectic (Ag) 1 (Cu) structure was observed after cooling. Furthermore, the liquidus temperature should have been increased above 900 °C such that its composition and temperature (marked as★) were located in the (L 1 Ag) region, as shown in Fig. 11(i). It was inferred that the clumpy TiCu4 was formed together with the Ag solid solution in the residual interlayer after cooling from 900 °C. As Ti was dissolved in the Ticusil melt, the Ti activity in the vicinity of the 3Y–ZrO2 interface also increased beyond a threshold value such that a thinner
layer of TiO was formed at 900 °C. The near Ag–Cu eutectic liquid (Lb) spread on the 3Y–ZrO2substrate by reactive wetting: xTidiss1 ZrO2! xTiO 1 ZrO2x, as shown in Fig. 11(b). Then, Cu atoms in the Ag–Cu liquid (L) reacted with Tidissto form a thinner layer of TiCu over the TiO layer, as observed in Fig. 11(c). Figure 11(d) demonstrates that the interfacial layer sequence of Ti2Cu/TiCu/Ti3Cu4/residual interlayer (Ag 1 TiCu4)/TiCu/TiO existed after brazing at 900 °C/0.1 h.
Figure 11(e) demonstrates that four continuous and thicker Ti–Cu layers (Ti2Cu/TiCu/Ti3Cu4/TiCu4) were formed sequentially from the Ti side to the interlayer upon brazing at 900 °C/1 h, which means that dramatic Cu consumption would cause complete solidification with the composition of the residual Ag–Cu alloy (L) being located in the a-Ag solid solution of the Ag–Cu phase diagram (marked by ▲).12 Neither the clumpy TiCu4 nor the dendritic TiCu4 was observed because no liquid phase remained at this stage. After cooling, supersaturated Cu and O were precipitated as Cu2O in the matrix of Ag. In addition to the TiO layer, two oxide layers of Ti2O and Ti3Cu3O were observed on the 3Y–ZrO2side, which indicated that significant O atoms (from the ZrO2x) diffused across the TiO layer to react with the Cu and Ti atoms of the residual interlayer such that the residual interlayer should be considered an Ag alloy with Ti, Cu, and O in solid solutions. The formation of Ti3Cu3O could be described by the following reaction: 3TiCu 1 Odiss ! Ti3Cu3O. Thermodynamic calculations of the Gibbs free energy (ΔG) for Ti3Cu3O (502.5 kJ/mol) confirmed that it was more stable than TiCu (14.7 kJ/mol) at 900 °C.21,32Finally, the interfacial layers [Ti2Cu/TiCu/Ti3Cu4/TiCu4/residual inter-layer (Ag 1 Cu2O)/Ti3Cu3O/Ti2O/TiO] were observed after brazing at 900 °C/1 h, as shown in Fig. 11(f ).
Figure 11(g) shows the microstructure expected before cooling for brazing at 900 °C/6 h. The TiCu phase grew at the expense of Ti3Cu4 and TiCu4 at the Ti/Ticusil in-terface, indicating that TiCu was much more stable than both Ti3Cu4and TiCu4.
21,33
While three interfacial oxide layers (Ti3Cu3O, Ti2O, and TiO) became thicker, no new phase was formed in the vicinity of the 3Y–ZrO2. However, a two-phase region ofa-Ti 1 Ti2Cu appeared on the Ti side due to a hypoeutectoid reaction during cooling, implying that the extensive diffusion of Cu occurred on the Ti side. The composition of the residual interlayer went further into the (Ag) region and was located at the position marked by a dot (●) in the Ag–Cu
phase diagram in Fig. 11(i). Fine Cu2O and Ti3Cu3O phases were precipitated in the Ag solid solution of the residual interlayer during cooling. Finally, the interfacial layers of a-Ti 1 Ti2Cu/Ti2Cu/TiCu/residual interlayer (Ag1 Cu2O1 Ti3Cu3O)/Ti3Cu3O/Ti2O/TiO in sequence were observed after brazing at 900 °C/6 h, as shown in Fig. 11(h).
FIG. 10. (a) A vertical Cu–Ti–(60% Ag) section of the Ag–Cu–Ti
system between 700 and 1300 °C and (b) the cooling paths (U4–U7) in
the partial liquidus projection of the Ag–Cu–Ti phase diagram between
779 and 860 °C. Different colors are used for the one-phase region (purple), two-phase region (white), and three-phase region (yellow).
IV. CONCLUSIONS
In this study, 3 mol% Y2O3-stabilized zirconia (3Y– ZrO2) and commercially pure titanium (cp-Ti) joints were fabricated with an Ag68.8Cu26.7Ti4.5interlayer at 900 °C for various brazing periods. After brazing at 900 °C/0.1 h, Ti2Cu, TiCu, Ti3Cu4, and TiCu4layers were present at the Ti/Ticusil interface, while TiCu and TiO layers were observed at the Ticusil/3Y–ZrO2 interface. After brazing at 900 °C/1 h, Ti3Cu3O and Ti2O layers were formed at the interlayer/ZrO2interface. After brazing at 900 °C/6 h, the TiCu layer grew at the expense of Ti3Cu4and TiCu4at the Ti/interlayer interface, and a two-phase (a-Ti 1 Ti2Cu) region was observed on the Ti side. While clumpy TiCu4 was formed in the melt interlayer, the liquid phase disappeared due to the extensive chemical reaction between Ti and Cu, leading to the formation of an Ag solid solution withfine Cu2O precipitates. The bonding mechanisms could be categorized into various stages. In stage 1, melting and liquid phase separation were the predominant mechanisms. In stage 2, dissolution and reactive wetting were the predominant mechanisms. In stage 3, interdiffusion and/or chemical reactions between Ti and Cu were the predominant mechanisms. In stage 4, solidification was the predominant mechanism after Ti and Cu were consumed due to the formation of various Ti–Cu intermetallic compounds, increasing the liquidus of residual Ag–Cu above the brazing temperature. In the cooling stage, fine Cu2O and/or Ti3Cu3O phases were precipitated in the
residual interlayer, while a eutectoid reaction (a-Ti 1 Ti2Cu) occurred on the Ti side for long-term brazing.
ACKNOWLEDGMENT
This research was supported by the National Science Council (Taiwan) under Contract No. NSC98-2221-E-009-039-MY2.
REFERENCES
1. O. Dezellus, J. Andrieux, F. Bosslet, M. Sacerdote-Peronnet,
T. Baffie, F. Hodaj, N. Eustathopoulos, and J.C. Viala: Transient
liquid phase bonding of titanium to aluminium nitride. Mater. Sci. Eng., A 495, 254 (2008).
2. R.K. Shiue, S.K. Wu, F.Y. Chen, and T.E. Yang: Interfacial reactions and wettability of 72Ag-28Cu braze on CP-Ti substrate using infrared heating. Metall. Mater. Trans. A 43, 1742 (2012). 3. C.C. Liu, C.L. Ou, and R.K. Shiue: The microstructural observation
and wettability study of brazing Ti-6Al-4V and 304 stainless steel using three braze alloys. J. Mater. Sci. 37(11), 2225 (2002).
4. D. Sotiropoulou and P. Nikolopoulos: Work of adhesion in ZrO2
liquid-metal systems. J. Mater. Sci. 28(2), 356 (1993).
5. W.D. Kingery, H.K. Bowen, and D.R. Uhlmann: Introduction to Ceramics, 2nd ed. (John Wiley & Sons, New York, 1991), p. 209. 6. A. Guedes, A.M.P. Pinto, M. Vieira, and F. Viana: Multulayered
interface in Ti/MacorÒmachinable glass-ceramic joints. Mater. Sci.
Eng., A 301, 118 (2001).
7. W.B. Hanson, K.I. Ironside, and J.A. Fernie: Active metal brazing
of zirconia. Acta. Mater. 48(18–19), 4673 (2000).
8. H.Y. Chan, D.W. Liaw, and R.K. Shiue: The microstructural observation of brazing Ti-6Al-4V and TZM using the BAg-8 braze alloy. Int. J. Refract. Met. Hard Mater. 22(1), 27 (2004).
FIG. 11. Microstructural evolution of the Ti/Ticusil/3Y–ZrO2joint at 900 °C for various brazing periods: (a)! (b) ! (c) ! (d) for 0.1 h, (a) ! (b)
! (e) ! (f) for 1 h, and (a) ! (b) ! (g) ! (h) for 6 h. The activity profiles of Ti and Cu are depicted on an arbitrary scale for simplicity. (i) the Ag-Cu phase diagram showing the corresponding compositions for various brazing periods.
9. Z.G. Wang, N. Kato, K. Sasaki, T. Hirayama, and H. Saka: Electron holographic mapping of two-dimensional doping areas in cross-sectional device specimens prepared by the lift-out technique based on a focused ion beam. J. Electron Microsc. 53(2), 115 (2004). 10. G. Cliff and G.W. Lorimer: The quantitative analysis of thin
specimens. J. Microsc. 103(2), 203 (1975).
11. J.I. Goldstein, D.B. Williams, and G. Cliff: Quantitative x-ray analysis. In Principles of Analytical Electron Microscopy, D.C. Joy, A.D. Romig, and J.I. Goldstein eds.; Plenum Press: New York, 1986; p. 155.
12. J.L. Murray: Calculations of stable and metastable equilibrium diagrams of the Ag-Cu and Cd-Zn systems. Metall. Trans. A 15(2), 261 (1984).
13. K.L. Lin, M. Singh, and R. Asthana: Interfacial characterization of YSZ-to-steel joints with Ag-Cu-Pd interlayers for solid oxide fuel cell applications. Ceram. Int. 38(3), 1991 (2012).
14. G.W. Liu, W. Li, G.J. Qiao, H.J. Wang, J.F. Yang, and T.J. Lu: Microstructures and interfacial behavior of zirconia/stainless steel joint prepared by pressureless active brazing. J. Alloys Compd. 470
(1–2), 163 (2009).
15. Y.W. Chang and C.C. Lin: Compositional dependence of phase formation mechanisms at the interface between titanium and calcia-stabilized zirconia at 1550°C. J. Am. Ceram. Soc. 93(11), 3893 (2010).
16. J.G. Lee, S.J. Hong, M.K. Lee, and C.K. Rhee: High strength bonding of titanium to stainless steel using an Ag interlayer. J. Nucl.
Mater. 395(1–3), 145 (2009).
17. J. Andrieux, O. Dezellus, F. Bosselet, M. Sacerdote-Peronnet, C. Sigala, R. Chiriac, and J.C. Viala: Details on the formation of
Ti2Cu3 in the Ag-Cu-Ti system in the temperature range
790-860°C. J. Phase Equilib. Diffus. 29(2), 156 (2008).
18. G.P. Kelkar, K.E. Spear, and A.H. Carim: Thermodynamic eval-uation of reaction products and layering in brazed alumina joints. J. Mater. Res. 9(9), 2244 (1994).
19. M.L. Santella, J.A. Horton, and J.J. Pak: Microstructure of alumina brazed with a silver-copper-titanium alloy. J. Am. Ceram. Soc. 73(6), 1785 (1990).
20. H.J. Lee and H.I. Aaronson: Eutectoid decomposition mechanisms in hypoeutectoid Ti-X alloys. J. Mater. Sci. 23(1), 150 (1988). 21. Y.H. Liang, H.Y. Wang, Y.F. Yang, Y.Y. Wang, and Q.C. Jiang:
Evolution process of the synthesis of TiC in the Cu-Ti-C system. J. Alloys Compd. 452(2), 298 (2008).
22. C. Hammerl, B. Renner, B. Rauschenbach, and W. Assmann: Phase
formation in titanium after high-fluence oxygen ion implantation.
Nucl. Instrum. Methods Phys. Res., Sect. B 148(1–4), 851 (1999).
23. B. Holmberg: Disorder and order in solid solutions of oxygen in a-titanium. Acta Chem. Scand. 16, 1245 (1962).
24. K.L. Lin and C.C. Lin: Reaction between titanium and zirconia powders during sintering at 1500°C. J. Am. Ceram. Soc. 90(7), 2220 (2007).
25. J.S. Pimenta, A.J.A. Buschinelli, R.M. do Nascimento,
A.E. Martinelli, and J. Remmel: Joining of zirconia mechanically metalized with titanium. Cerâmica 56, 212 (2010).
26. A. Soon, M. Todorova, B. Delley, and C. Stampfl: Thermodynamic
stability and structure of copper oxide surfaces: Afirst-principles
investigation. Phys. Rev. B 75, 125420 (2007).
27. V.N. Eremenko, Y.I. Buyanov, and N.M. Panchenko: Polythermal and isothermal sections of the system titanium-copper-silver. Part II. Sov. Powder Metall. Met. Ceram. 9(5), 410 (1970).
28. V.N. Eremenko, Y.I. Buyanov, and N.M. Panchenko: The liquidus surface of the system titanium-copper-silver. Powder Metall. Met. Ceram. 9(4), 301 (1970).
29. J.L. Murray and K.J. Bhansali: The Ag-Ti (silver-titanium) system. J. Phase Equil. 4(2), 178 (1983).
30. J.L. Murray: The Cu-Ti (copper-titanium) system. J. Phase Equil. 4(1), 81 (1983).
31. S. Hirnyj and J.E. Indacochea: Phase transformations in
Ag70.5Cu26.5Ti3 filler alloy during brazing processes. Chem.
Met. Alloys 1, 323 (2008).
32. M. Yang, T. Lin, and P. He: Microstructure evolution of Al2O3/
Al2O3joint brazed with Ag-Cu-Ti1 B 1 TiH2compositefiller.
Ceram. Int. 38, 289 (2012).
33. R.H. Shiue and S.K. Wu: Infrared brazing Ti50Ni50and Ti-6Al-4V
