Active Brazing of Alumina and Copper with Multicomponent Ag-Cu-Sn-Zr-Ti Filler

: The study was designed to investigate the synergic effect of Ti and Sn in the active metal brazing of Al 2 O 3 ceramic to copper brazed, using the multicomponent Ag-Cu-Zr ﬁller alloy. Numer-ous ﬁne and hexagonal-shaped rod-like ternary intermetallic (Zr, Ti) 5 Sn 3 phase (L/D = 5.1 ± 0.8, measured in microns) were found dispersed in the Ag-Cu matrix of Ag-18Cu-6Sn-3Zr-1Ti alloy, along with the ternary CuZrSn intermetallic phases. An approximate 15 ◦ reduction in contact angle and 3.1 ◦ C reduction in melting point are observed upon the incorporation of Ti and Sn in Ag-18Cu-3Zr ﬁller. Interestingly, the interface microstructure of Al 2 O 3 /Cu joints brazed by using Ag-18Cu-6Sn-3Zr-1Ti ﬁller shows a double reaction layer: a discontinuous Ti-rich layer consisting of (Cu, Al) 3 (Ti, Zr) 3 O, TiO, and in-situ Cu-(Ti, Zr) precipitates on the Al 2 O 3 side and continuous Zr-rich layer consisting of ZrO 2 on the ﬁller side. The shear strength achieved in Al 2 O 3 /Cu joints brazed with Ag-18Cu-6Sn-3Zr-1Ti ﬁller is 31% higher, compared to the joints brazed with Ag-18Cu-6Sn-3Zr ﬁller. Failure analysis reveals a composite fracture mode indicating a strong interface bonding in Al 2 O 3 /Ag-18Cu-6Sn-3Zr-1Ti ﬁller/Cu joints. The ﬁndings will be helpful towards the development of high entropy brazing ﬁllers in the future.


Introduction
Ceramics have a wide range of applications in electric vehicles, aerospace industries, nuclear and chemical-powder plant applications. Irrespective of their potential in various applications, utilizing ceramics in required complex shapes imposes serious manufacturing and economic limitations in industrial sectors [1][2][3][4][5]. The idea of ceramic to metal joining was considered as a promising approach that resolves the existing difficulties in the manufacturing process. Joining metals to ceramics provides complementary properties that cannot be achieved with either ceramics or metal alone; however, achieving a good joint interface between ceramic and metal has a lot of hurdles to overcome [2,3]. Moreover, the overall properties rely on the joint interface strength. Ceramics and metals have distinct differences in their bonding nature and chemical and physical properties. Metals have a metallic bond with free electrons; in contrast, ceramics have a covalent or ionic bond with extremely stable outer-shell electrons. Hence, molten metal does not wet the ceramic surface, thus limiting the prerequisite condition for joining, i.e., establishing intimate contact between ceramic and metal surface [4]. Additionally, distinct physical properties between ceramics and metals, such as elastic modulus, strength, and Coefficient of Thermal Expansion (CTE) mismatch lead to large residual stress during cooling, affecting the joint strength [5].
Adhesive bonding [6], diffusion bonding [7], active metal brazing (AMB) [8] and transient liquid-phase bonding [9] are some of the well-known ceramic-to-metal joining methods. Owing to its simplicity and good joint properties, AMB is the commonly employed ceramic-metal joining technique. AMB is a joining process wherein ceramic and mechanically mixed and melted at 1000 • C, in a vacuum furnace, at 1.3 × 10 −2 Pa. Vacuum melting is necessary to prevent the oxidation of filler alloys. The alloys were homogenized and re-melted at 1000 • C, to achieve compositional uniformity. From the ingot, the braze sheets were then cut, polished into sheet of dimension 15 mm × 15 mm × 0.3 mm and ultrasonically cleaned in ethanol. Then, 96% pure Al 2 O 3 (purchased from Korea electronic material, Gyeonggi, Korea), having a dimension of 15 mm × 15 mm × 5 mm, and 99.9% pure Cu (purchased from Daeduck metals, Gyeonggi, Korea) with a dimension of 20 mm × 20 mm × 5 mm were used as base ceramic and metal substrates respectively. These substrates were polished by using #500-, #800-and #1000-grit SiC abrasive paper and ultrasonically cleaned in ethanol. The as-prepared braze sheets were then sandwiched between Al 2 O 3 and Cu substrates. AMB was performed at the vacuum level of 1.3 × 10 −3 Pa with the heating and cooling rate of 10 and 5 • C/min, respectively. The peak brazing temperature and brazing time were fixed as 840 • C and 15 min, respectively, for both the brazing alloys.

Microstructural and Thermal Analysis
The brazed samples were cut perpendicular to the bonded interface, and metallographic specimens were prepared. Microstructures of the filler alloy and brazed joints were examined, using analytical SEM (JEOL JSM-6010PLUS, Tokyo, Japan) attached with an energy-dispersive spectrometer (EDS). The microstructure specimens were etched in a solution containing 25 mL H 2 O + 50 mL H 2 O 2 + 25 mL ammonia water prior to the microstructural examination. The reaction phase at the Al 2 O 3 /filler interface was identified by XRD, using the Cu K-α spectrum having a wavelength of 1.5406 A (Bruker AXS GmbH-D8 DISCOVER, Karlsruhe, Germany). For XRD, the brazed samples were cut parallel to the bonded interface and polished until the reaction phases at the interface were revealed. The melting point of the filler alloys was analyzed, using differential scanning calorimetry (DSC, SDT 650, TA instruments, New castle, DE, USA) at a heating rate and cooling rate of 10 • C/min under N 2 atmosphere. The sample weight was maintained at 10.5 ± 0.2 mg in all the DSC tests.

Contact Angle Measurement and Shear Test
The spreading tests for contact angle measurement were carried in accordance with the Japanese Industrial Standard (2003: JIS-Z-3198-3). To measure the contact angle, spreading of 0.3 g of filler alloy on Al 2 O 3 substrate was carried in a vacuum furnace, at a vacuum level 1.3 × 10 −3 Pa, with a heating and cooling rate of 10 and 5 • C/min, respectively. The temperature and time were 840 • C and 15 min, respectively. After spreading, the samples were mounted in epoxy, cut perpendicular to the spreading interface, and the contact angle is measured by following Young's equation: [22] where γ SL , γ SV and γ LV represent the solid-liquid, solid-vapor and liquid-vapor surface tension, respectively, and θ represents the contact angle. The measurement procedure is schematically illustrated in Figure 1. Assuming the solidified solder having a spherical shape, a best fit circle for solder surface was drawn. A tangent was drawn at the point where R is the radius of the circle and y c is the y-axis coordinate of center of the circle. procedure is schematically illustrated in Figure 1. Assuming the solidified solder having a spherical shape, a best fit circle for solder surface was drawn. A tangent was drawn at the point P, which is the intersection of the circle and the top surface of the Al2O3 substrate. The contact angle is then given by Equation (2) [23]: where R is the radius of the circle and yc is the y-axis coordinate of center of the circle. The reported contact angle measurement was the average of three specimens. The shear test was carried at a shear speed of 0.5 mm/min using UTM WDW-20 (Chenda tester, Jian, China). Each reported shear test value is an average value of three tested specimens. After the shear test, the fractured Al2O3 samples were mounted and sectioned to study the crack propagation.

Results and Discussion
3.1. Microstructure of Ag-18Cu-6Sn-3Zr and Ag-18Cu-6Sn-3Zr-1Ti Brazing Filler Figure 2 shows the microstructure and EDS analysis of as-cast filler alloys. The microstructure of the Ag-18Cu-6Sn-3Zr filler alloy as shown in Figure 2 consists of a bright Ag-rich and dark Cu-rich matrix region. It reveals a bright lamellar phase marked as A and dark blocky phase marked as B dispersed In the Ag-Cu matrix. EDS analysis on points A and B are shown in Figure 2. Based on the at.% calculations from the EDS analysis and the references listed in Table 2, the bright lamellar phase indicated by point A refers to the ternary CuZrSn phase [23] and the blocky phase indicated by point B refers to a ternary Cu4AgZr IMC [24]. The reported contact angle measurement was the average of three specimens. The shear test was carried at a shear speed of 0.5 mm/min using UTM WDW-20 (Chenda tester, Jian, China). Each reported shear test value is an average value of three tested specimens. After the shear test, the fractured Al 2 O 3 samples were mounted and sectioned to study the crack propagation.

Results and Discussion
3.1. Microstructure of Ag-18Cu-6Sn-3Zr and Ag-18Cu-6Sn-3Zr-1Ti Brazing Filler Figure 2 shows the microstructure and EDS analysis of as-cast filler alloys. The microstructure of the Ag-18Cu-6Sn-3Zr filler alloy as shown in Figure 2 consists of a bright Ag-rich and dark Cu-rich matrix region. It reveals a bright lamellar phase marked as A and dark blocky phase marked as B dispersed In the Ag-Cu matrix. EDS analysis on points A and B are shown in Figure 2. Based on the at.% calculations from the EDS analysis and the references listed in Table 2, the bright lamellar phase indicated by point A refers to the ternary CuZrSn phase [23] and the blocky phase indicated by point B refers to a ternary Cu 4 AgZr IMC [24]. Ag-Cu system is an interesting combination having the same crystal structure (facecentered cubic) and positive enthalpy of mixing (10.8 KJ·mol −1 [23]). During solidification, to attain equilibrium Cu and Ag atoms tend to segregate and forms Cu-rich and Ag-rich matrix [25]. As there are no thermodynamic investigations available for a quaternary Ag-Cu-Sn-Zr alloy, the present work characterizes the intermetallic phases based on the ternary systems of Ag-Cu-Zr [26][27][28] and Cu-Sn-Zr [24,29]. Zr has a maximum solid solubility of 0.1 at.% at 900 • C in Ag [26] and 0.12 at.% at 972 • C in Cu [27]. Liquid-phase separation is observed in the ternary Ag-Cu-Zr system in Ag-rich area, where the liquid melt (L) is separated into Ag-rich liquid (L 1 ) and Cu-Zr-rich liquid (L 2 ) [28]. During solidification, precipitation of Cu 4 AgZr IMC and occurs at 898 • C by the reaction (3). Cu 4 AgZr is a high-temperature IMC with a congruent melting temperature of 1080 • C [22]. Ag-Cu system is an interesting combination having the same crystal structure (facecentered cubic) and positive enthalpy of mixing (10.8 KJ·mol −1 [23]). During solidification, to attain equilibrium Cu and Ag atoms tend to segregate and forms Cu-rich and Ag-rich matrix [25]. As there are no thermodynamic investigations available for a quaternary Ag-Cu-Sn-Zr alloy, the present work characterizes the intermetallic phases based on the ternary systems of Ag-Cu-Zr [26][27][28] and Cu-Sn-Zr [24,29]. Zr has a maximum solid solubility of 0.1 at.% at 900 °C in Ag [26] and 0.12 at.% at 972 °C in Cu [27]. Liquid-phase separation is observed in the ternary Ag-Cu-Zr system in Ag-rich area, where the liquid melt (L) is separated into Ag-rich liquid (L1) and Cu-Zr-rich liquid (L2) [28]. During solidification, precipitation of Cu4AgZr IMC and occurs at 898 °C by the reaction (3). Cu4AgZr is a high-temperature IMC with a congruent melting temperature of 1080 °C [22]. L (Ag) rich L (Cu-Zr) rich + Cu AgZr + α-Ag  Romaka et al. [29] experimentally investigated the isothermal section of the ternary Cu-Sn-Zr system at 397 and 497 °C . Apparently, from the investigation, a ternary equiatomic CuZrSn phase (designed as τ1) with TiNiSi structure type and cell parameters of a = 0.66279 nm, b = 0.43679 nm, c = 0.76791 nm has been identified. Yuan et al. [24] experimentally observed the presence of CuZrSn (τ1) phase in the Cu-Sn-Zr system at 700 °C . The atomic percentages of the CuZrSn (τ1) phase as reported by Romaka et al. [29] and Yuan et al. [24] is found to be good in agreement with the EDS in the present work as shown in Table 2. From the microstructure, the lamellar CuZrSn (τ1) phase either exists along with the Cu-rich inclusions or formed from Cu4AgZr IMC. The CuZrSn (τ1) phase could have transformed from the Sn-enriched Cu-Zr liquid (L2) during solidification. However, an in-depth analysis is required to understand the transformation mechanism. Figure 3a-h shows the SEM microstructure, elemental mapping and EDS analysis of as-cast Ag-18Cu-6Sn-3Zr-1Ti filler alloy. Similar to Ag-18Cu-6Sn-3Zr filler, CuZrSn (τ1) phase (marked as C) is observed in Ag-18Cu-6Sn-3Zr1Ti filler as shown in Figure 3b. Since Ti and Zr are soluble over the entire composition range, presence of Ti in CuZrSn (τ1) phase is inevitable as seen from the EDS analysis of points C in Figure 3h. Blocky Cu4AgZr phase is not seen in Ag-18Cu-6Sn-3Zr-1Ti filler. Instead, a fine and hexagonal rod-like phase (L/D = 5.1 ± 0.8 measured in microns) was observed. EDS analysis indicated by point D confirms the presence of Ti, Zr and Sn in the rod-like phase. The at.% of the elements indicates that this phase is identical to (Zr, Ti)5Sn3 phase composition in the literature, as shown in Table 2 [30,31]. Elemental mapping of the magnified microstructure as shown in Figure 3c-g confirms the coexistence of Zr, Ti and Sn in (Zr, Ti)5Sn3 phase. Nowotny et al. reported that Zr5Sn3 and Ti5Sn3 IMCs can substitute Ti and Zr atoms to form a hexagonal (Zr, Ti)5Sn3 phase [31]. Several other researchers also support the formation of (Zr, Ti)5Sn3 phase in the Sn-Zr-Ti system over a wide range of Ti and Zr compositions [31][32][33]. Fine (Zr, Ti)5Sn3 phase formed upon the addition of Ti and Sn in Ag-Cu-Zr filler could Romaka et al. [29] experimentally investigated the isothermal section of the ternary Cu-Sn-Zr system at 397 and 497 • C. Apparently, from the investigation, a ternary equiatomic CuZrSn phase (designed as τ 1 ) with TiNiSi structure type and cell parameters of a = 0.66279 nm, b = 0.43679 nm, c = 0.76791 nm has been identified. Yuan et al. [24] experimentally observed the presence of CuZrSn (τ 1 ) phase in the Cu-Sn-Zr system at 700 • C. The atomic percentages of the CuZrSn (τ 1 ) phase as reported by Romaka et al. [29] and Yuan et al. [24] is found to be good in agreement with the EDS in the present work as shown in Table 2. From the microstructure, the lamellar CuZrSn (τ 1 ) phase either exists along with the Cu-rich inclusions or formed from Cu 4 AgZr IMC. The CuZrSn (τ 1 ) phase could have transformed from the Sn-enriched Cu-Zr liquid (L 2 ) during solidification. However, an in-depth analysis is required to understand the transformation mechanism. Figure 3a-h shows the SEM microstructure, elemental mapping and EDS analysis of as-cast Ag-18Cu-6Sn-3Zr-1Ti filler alloy. Similar to Ag-18Cu-6Sn-3Zr filler, CuZrSn (τ 1 ) phase (marked as C) is observed in Ag-18Cu-6Sn-3Zr1Ti filler as shown in Figure 3b. Since Ti and Zr are soluble over the entire composition range, presence of Ti in CuZrSn (τ 1 ) phase is inevitable as seen from the EDS analysis of points C in Figure 3h. Blocky Cu 4 AgZr phase is not seen in Ag-18Cu-6Sn-3Zr-1Ti filler. Instead, a fine and hexagonal rod-like phase (L/D = 5.1 ± 0.8 measured in microns) was observed. EDS analysis indicated by point D confirms the presence of Ti, Zr and Sn in the rod-like phase. The at.% of the elements indicates that this phase is identical to (Zr, Ti) 5 Sn 3 phase composition in the literature, as shown in Table 2 [30,31]. Elemental mapping of the magnified microstructure as shown in Figure 3c-g confirms the coexistence of Zr, Ti and Sn in (Zr, Ti) 5 Sn 3 phase. Nowotny et al. reported that Zr 5 Sn 3 and Ti 5 Sn 3 IMCs can substitute Ti and Zr atoms to form a hexagonal (Zr, Ti) 5 Sn 3 phase [31]. Several other researchers also support the formation of (Zr, Ti) 5 Sn 3 phase in the Sn-Zr-Ti system over a wide range of Ti and Zr compositions [31][32][33]. Fine (Zr, Ti) 5 Sn 3 phase formed upon the addition of Ti and Sn in Ag-Cu-Zr filler could be more beneficial in mechanical properties as compared to the blocky Cu 4 AgZr IMC in Ag-Cu-Zr and Ag-Cu-Zr-Ti fillers [22,34]. be more beneficial in mechanical properties as compared to the blocky Cu4AgZr IMC in Ag-Cu-Zr and Ag-Cu-Zr-Ti fillers [22,34].

Thermal analysis and Contact
Angle of Ag-18Cu-6Sn-3Zr and Ag-18Cu-6Sn-3Zr-1Ti Filler on Al2O3 Substrate Figure 4a represents the DSC curves for Ag-18Cu-6Sn-3Zr and Ag-18Cu-6Sn-3Zr-1Ti filler. One endothermic peak is observed in both the filler alloys. The peak melting temperature of Ag-18Cu-6Sn-3Zr filler is 782.9 °C slightly higher than the Ag-Cu eutectic melting temperature of 779 °C. Addition of 1 wt.% Ti in Ag-18Cu-6Sn-3Zr filler has slightly reduced the peak melting temperature to 779.8 °C . Reactive wetting and spreading are considered an important process in ceramic-metal brazing. Though distinctly defined, spreading and wetting are interrelated concepts and are often interpreted with a certain degree of ambiguity. Wetting is a necessary precursory condition required for spreading [35]. Landry proposed a model wherein the formation of a new compound at the interface leads to wetting in metal/ceramic systems. The final contact angle and spreading kinetics depend on the newly formed interfacial compound [36]. The lower contact angle is a consequence of a higher spreading rate and wettability. Brazing atmosphere has a substantial influence on the wettability of filler metal on the ceramic surface. Low oxygen potential in brazing atmosphere can prevent the active elements (Ti and Zr) in the filler from reacting with the oxygen, which may result in a loss in the wettability [22]. Figure 4b shows the contact angle of Ag-18Cu-6Sn-3Zr and Ag-

Thermal Analysis and Contact
Angle of Ag-18Cu-6Sn-3Zr and Ag-18Cu-6Sn-3Zr-1Ti Filler on Al 2 O 3 Substrate Figure 4a represents the DSC curves for Ag-18Cu-6Sn-3Zr and Ag-18Cu-6Sn-3Zr-1Ti filler. One endothermic peak is observed in both the filler alloys. The peak melting temperature of Ag-18Cu-6Sn-3Zr filler is 782.9 • C slightly higher than the Ag-Cu eutectic melting temperature of 779 • C. Addition of 1 wt.% Ti in Ag-18Cu-6Sn-3Zr filler has slightly reduced the peak melting temperature to 779.8 • C. Reactive wetting and spreading are considered an important process in ceramic-metal brazing. Though distinctly defined, spreading and wetting are interrelated concepts and are often interpreted with a certain degree of ambiguity. Wetting is a necessary precursory condition required for spreading [35]. Landry proposed a model wherein the formation of a new compound at the interface leads to wetting in metal/ceramic systems. The final contact angle and spreading kinetics depend on the newly formed interfacial compound [36]. The lower contact angle is a consequence of a higher spreading rate and wettability. Brazing atmosphere has a substantial influence on the wettability of filler metal on the ceramic surface. Low oxygen potential in brazing atmosphere can prevent the active elements (Ti and Zr) in the filler from reacting with the oxygen, which may result in a loss in the wettability [22]. Figure 4b shows the contact angle of Ag-18Cu-6Sn-3Zr and Ag-18Cu-6Sn-3Zr-1Ti filler on the Al 2 O 3 substrate, at 840 • C, for 15 min. A contact angle of 70.1 ± 0.4 • was obtained for the Ag-18Cu-6Sn-3Zr filler. This value is considerably higher than the contact angle of 54.2 ± 0.5 • observed for Ag-18Cu-6Sn-3Zr-1Ti filler. In metal-ceramic systems, the spreading rate is limited by the rate of interfacial reaction process namely: diffusive transport of reacting species and local reaction kinetics at the interface [36]. Although Ti and Zr are both active elements, low atomic mass and small atomic radius of Ti assist in rapid diffusion through the Ag-Cu alloy [17]. 18Cu-6Sn-3Zr-1Ti filler on the Al2O3 substrate, at 840 °C , for 15 min. A contact angle of 70.1 ± 0.4° was obtained for the Ag-18Cu-6Sn-3Zr filler. This value is considerably higher than the contact angle of 54.2 ± 0.5° observed for Ag-18Cu-6Sn-3Zr-1Ti filler. In metalceramic systems, the spreading rate is limited by the rate of interfacial reaction process namely: diffusive transport of reacting species and local reaction kinetics at the interface [36]. Although Ti and Zr are both active elements, low atomic mass and small atomic radius of Ti assist in rapid diffusion through the Ag-Cu alloy [17]. The activity coefficient is an important parameter that quantifies the activity (χTi) of titanium at the filler alloy and substrate interface during brazing and is given by Equations (4) and (5) [37]:

Ti
Ti Ti χ = γ .X (4) where XTi is the concentration of Ti and Ti γ is the activity coefficient described as follows [17].  The activity coefficient is an important parameter that quantifies the activity (χ Ti ) of titanium at the filler alloy and substrate interface during brazing and is given by Equations (4) and (5) [37]: where X Ti is the concentration of Ti and γ Ti is the activity coefficient described as follows [17].
where G i E is the excess Gibbs free energy of mixing, R is the gas constant and T is the temperature. The higher the activity coefficient, the more there will be segregation of active elements at the ceramic-metal interface. The activity coefficient of Ti in infinitely diluted Cu-Ti alloy is −1.4 [38], whereas the activity coefficient of Zr in infinitely diluted Cu-Zr alloy is −4 [39]. Due to its high activity, Ti facilitates the formation of reaction products on the Al 2 O 3 substrates at a faster rate, as compared to Zr. Consequently, the spreading enhances reducing the contact angle in Ag-18Cu-6Sn-3Zr-1Ti filler.
3.3. Interface Analysis of Al 2 O 3 /Cu Joint 3.3.1. Interfacial Study of Cu/Al 2 O 3 Joints Brazed with Ag-18Cu-6Sn-3Zr Filler Figure 5 shows the characteristic features of Al 2 O 3 /Cu joint brazed with Ag-18Cu-6Sn-3Zr filler alloy. The chemical reactions between Al 2 O 3 and Ag-18Cu-6Sn-3Zr filler resulted in a reaction layer with a thickness of 5.2 ± 0.4 µm at the Al 2 O 3 /Cu interface, as shown in Figure 5a. XRD analysis of reaction zone shown in Figure 5b displays the peaks corresponding to α-Al 2 O 3 , Ag, Cu and monoclinic ZrO 2 . Moreover, α-Al 2 O 3 peak originates from the substrate, while Ag and Cu peaks originated from the filler alloy. EDS analysis of the interface for the individual elements Al, O, Zr, Cu, Ag and Sn in Figure 5c, shows the presence of Zr and O in the ZrO 2 reaction layer. A copper-rich region is also observed in the immediate vicinity of the Al 2 O 3 substrate ahead of ZrO 2 layer. Formation of ZrO 2 reaction layer at Al 2 O 3 /filler interface is considered as more complex when compared to a simple Al 2 O 3 reduction-Zr oxidation reaction [15]. During brazing, Zr in the molten filler diffuses towards Al 2 O 3 substrate, facilitates wetting and penetrates along the grain boundaries of Al 2 O 3 . Meanwhile, Al 2 O 3 gets reduced to Al and O ions by Reaction (6) [40]: Zr in the molten filler reacts with O ions and gets oxidized to ZrO 2 by Reaction (7): This leads to the oxidation-reduction reaction at the interface described below: However, Al 2 O 3 is thermodynamically more stable than ZrO 2 and the Gibbs free energy (∆G 0 R ) of Equation (8) is positive (396.8 kJ at 1127 • C) [15]. Loehman et al. proposed that the formation of Cu-Zr-Ag-Al precipitates on Al 2 O 3 can lower the Gibbs free energy of Equation (7) and favors the formation of ZrO 2 at the interface [15]. Therefore, after the filler metal is melted, Cu-Zr-rich liquid infiltrates along the Al 2 O 3 grain boundaries and the high oxidation potential of Zr could have led to redox reaction within Al 2 O 3 . Thus, the ZrO 2 reaction layer form and continue to grow by the outward diffusion of oxygen ions, leaving behind Cu at the Al 2 O/reaction layer interface to form Cu-rich precipitates [40]. Although a possible mechanism of ZrO 2 formation is supported based on Equation (8), sufficient thermodynamic evidence on the formation of ZrO 2 on Al 2 O 3 is still too complex to comprehend.

Interfacial Study of Cu/Al2O3
Joint Brazed with Ag-18Cu-6Sn-3Zr-1Ti Filler Figure 6 shows the characteristic features of Al2O3/Cu joint brazed with Ag-18Cu-6Sn-3Zr-1Ti filler. The microstructure of the Al2O3/Cu joint, as shown in Figure 6a, displays a reaction layer of thickness 7.3 ± 0.3 µ m at the interface. This is ~1.9 µ m thicker, as compared to the reaction layer obtained by using Ag-18Cu-6Sn-3Zr filler. Furthermore, unlike the joints brazed by using Ag-18Cu-6Sn-3Zr filler, no blocky IMCs are found in the brazed seam. XRD analysis of the reaction layer shown in Figure 6b displays the peaks of ɑ-Al2O3, Ag, Cu, Cu3Ti3O, CuTi3, TiO and m-ZrO2. High-magnification FESEM images and the elemental mapping, as shown in Figure 6c,d show two distinct composition of reaction

Interfacial Study of Cu/Al 2 O 3 Joint Brazed with
Ag-18Cu-6Sn-3Zr-1Ti Filler Figure 6 shows the characteristic features of Al 2 O 3 /Cu joint brazed with Ag-18Cu-6Sn-3Zr-1Ti filler. The microstructure of the Al 2 O 3 /Cu joint, as shown in Figure 6a, displays a reaction layer of thickness 7.3 ± 0.3 µm at the interface. This is~1.9 µm thicker, as compared to the reaction layer obtained by using Ag-18Cu-6Sn-3Zr filler. Furthermore, unlike the joints brazed by using Ag-18Cu-6Sn-3Zr filler, no blocky IMCs are found in the brazed seam. XRD analysis of the reaction layer shown in Figure 6b displays the peaks of α-Al 2 O 3 , Ag, Cu, Cu 3 Ti 3 O, CuTi 3 , TiO and m-ZrO 2 . High-magnification FESEM images and the elemental mapping, as shown in Figure 6c,d show two distinct composition of reaction layers (marked as I and II). A discontinuous Ti-rich region I is present immediately adjacent to Al 2 O 3 substrate, followed by a continuous Zr-rich region II. Figure 6d displays the EDS analysis of the phases in the reaction layer of the magnified image. The atomic percentage of elements from EDS analysis corresponding to the spots E, F and G are given in Table 3. The stoichiometric ratio of elements Zr, Ti, Cu, Al and O in spot E corresponds to Cu 3 Ti 3 O, where Ti is substituted by Zr, and Cu is substituted by Al. During reduction of Al 2 O 3 , Al ions diffuses towards the filler and participates in the reaction products formed during the brazing. Earlier research also confirms the solubility of Al in (Cu, Al) 3 Ti 3 O compound [41]. The presence of elements Cu, Zr, Ti and Al in the EDS analysis of spot F and their stoichiometric composition matches with the Cu-(Ti, Zr) IMC with Al substituted in Cu. The EDS analysis in the spot G indicates the presence of Zr and O that coincides with the stoichiometric percentage of ZrO 2 . By combining EDS, XRD and elemental mapping analysis, it is found that the phases present in region I could be (Cu, Al) 3 (Zr, Ti) 3 O, TiO and in situ Cu-(Ti, Zr) IMC and the region II mainly composed of ZrO 2 .
in Table 3. The stoichiometric ratio of elements Zr, Ti, Cu, Al and O in spot E corresponds to Cu3Ti3O, where Ti is substituted by Zr, and Cu is substituted by Al. During reduction of Al2O3, Al ions diffuses towards the filler and participates in the reaction products formed during the brazing. Earlier research also confirms the solubility of Al in (Cu, Al)3Ti3O compound [41]. The presence of elements Cu, Zr, Ti and Al in the EDS analysis of spot F and their stoichiometric composition matches with the Cu-(Ti, Zr) IMC with Al substituted in Cu. The EDS analysis in the spot G indicates the presence of Zr and O that coincides with the stoichiometric percentage of ZrO2. By combining EDS, XRD and elemental mapping analysis, it is found that the phases present in region I could be (Cu, Al)3(Zr, Ti)3O, TiO and in situ Cu-(Ti, Zr) IMC and the region II mainly composed of ZrO2.  As explained in Section 3.2, a high activity coefficient, low atomic mass and radius of Ti enable faster diffusion towards Al2O3 substrate, compared to Zr. Segregation of Ti results in the formation of Cu3Ti3O and TiO oxides in the immediate vicinity of Al2O3. Apparently, many works in the literature support the formation of interfacial reaction products, such as TiO2, TiO, Ti2O and Cu3Ti3O during the AMB between alumina and Ag-   [37,[42][43][44]. The established mechanism of reaction layer in Al 2 O 3 /Ag-Cu-Ti system is the formation of nm-thick TiO layer followed by µm-thick Ti 3 (Cu+Al) 3 O on Al 2 O 3 [43]. However, the formation mechanism can be affected by the composition of the filler alloy and joining conditions. After melting, Ti in the filler metal diffuses and reacts with alumina to form TiO. Upon cooling, TiO reacts with Cu-Ti intermetallic compounds to form Cu 3 Ti 3 O and Ti 2 O as expressed by the Equations (9) and (10) [44] As explained in Section 3.2, with low atomic mass and radius, Ti diffuses towards the Al 2 O 3 at a faster rate, and the high activity coefficient of Ti facilitates the formation of reaction products, such as TiO and Cu 3 Ti 3 O, on the Al 2 O 3 substrates. Hence, region I in the immediate vicinity of Al 2 O 3 substrate is Ti-rich. However, it is complex to understand the formation mechanism of in situ Cu-(Ti, Zr) IMC. With higher atomic mass and radius, Zr diffuses towards the Al 2 O 3 at a smaller rate and establish the formation of ZrO 2 in region II. This results in the formation of double reaction layers at the interface. Figure 7a displays the shear strength of Al 2 O 3 /Cu joints brazed at 840 • C for 15 min, using Ag-Cu-Sn-Zr filler and Ag-Cu-Sn-Zr-Ti filler as a function of Zr content. The average shear strength of Al 2 O 3 /Cu joints brazed with Ag-18Cu-6Sn-3Zr-1Ti filler is higher (18.4 MPa), compared with Cu/Al 2 O 3 joints brazed with Ag-18Cu-6Sn-3Zr filler (14.1 MPa). The reliability of the Al 2 O 3 /Cu joints depends on the residual stresses due to CTE and elastic modulus mismatch developed during cooling [5]. Crack propagation in the fractured Al 2 O 3 /Cu joints brazed with Ag-18Cu-6Sn-3Zr and Ag-18Cu-6Sn-3Zr-1Ti filler are shown in Figure 7b,c, respectively. In the specimens brazed with Ag-18Cu-6Sn-3Zr filler, cracks having initiated at the ZrO 2 reaction layer/filler interface propagated through the ZrO 2 layer, indicating a weak interface. Besides this, cracks are also seen in Al 2 O 3 ceramic substrate adjacent to the reaction layer.

Effect of Ag-Cu-Sn-Zr-Ti Filler on the Shear Strength of Al 2 O 3 /Cu Joints
MPa). The reliability of the Al2O3/Cu joints depends on the residual stresses due to CTE and elastic modulus mismatch developed during cooling [5]. Crack propagation in the fractured Al2O3/Cu joints brazed with Ag-18Cu-6Sn-3Zr and Ag-18Cu-6Sn-3Zr-1Ti filler are shown in Figure 7b,c, respectively. In the specimens brazed with Ag-18Cu-6Sn-3Zr filler, cracks having initiated at the ZrO2 reaction layer/filler interface propagated through the ZrO2 layer, indicating a weak interface. Besides this, cracks are also seen in Al2O3 ceramic substrate adjacent to the reaction layer. ZrO2 exhibits three polymorphs: monoclinic (m-ZrO2), tetragon (c-ZrO2). Moreover, m-ZrO2 is a stable structure at room temper During cooling in AMB, polymorphic transformation of t-ZrO2 to mby a volume increase of approximately 3%-5% [41]. As a consequenc residual stresses can be generated to this volume change and can re formation in the m-ZrO2 reaction layer [5]. Therefore, the fracture t becomes weak and is highly susceptible to crack propagation. Moreo and CTE mismatch can together generate large residual stresses tha Al2O3 ceramic substrate ahead of the reaction layer. Additionally, pre in the filler must have created plastic strain localization at the IMC/m In the specimens brazed with Ag-18Cu-6Sn-3Zr-1Ti filler, the ZrO2/filler interface, primarily propagated in ZrO2 reaction layer and towards the Al2O3 substrate. This composite fracture mode implie strong bonding at the interface. The CTE values of Al2O3, ZrO2 and 10 −6 /°C [43], 10.5 × 10 −6 / °C [10] and (19.2-22) × 10 −6 / °C [43], respectiv 18Cu-6Sn-3Zr-1Ti filler interface, the presence of the Ti-rich region c (15.2 × 10 −6 /°C [43]), TiO (9.1 × 10 −6 /°C [43]) oxides and Cu-(Ti, Zr) ah layer may have alleviated the residual stresses and micro-cracks in m generated due to the CTE mismatch and t-ZrO2 to m-ZrO2 transform Additionally, in situ formed Cu-(Ti, Zr) IMC dispersed in the oxid block the crack propagation in the reaction layer. This results in a s high fracture toughness. Thus, multicomponent Ag-Cu-Sn-Zr-T potential to explore the possibilities of upgrading the existing comm wide range of ceramics and metals.

Conclusions
Al2O3/Cu joint was successfully brazed at 840 °C , for 15 min, us ZrO 2 exhibits three polymorphs: monoclinic (m-ZrO 2 ), tetragonal (t-ZrO 2 ) and cubic (c-ZrO 2 ). Moreover, m-ZrO 2 is a stable structure at room temperature and pressure. During cooling in AMB, polymorphic transformation of t-ZrO 2 to m-ZrO 2 is accompanied by a volume increase of approximately 3%-5% [41]. As a consequence of volume change, residual stresses can be generated to this volume change and can result in micro-cracks formation in the m-ZrO 2 reaction layer [5]. Therefore, the fracture toughness of m-ZrO 2 becomes weak and is highly susceptible to crack propagation. Moreover, volume change and CTE mismatch can together generate large residual stresses that further weaken the Al 2 O 3 ceramic substrate ahead of the reaction layer. Additionally, presence of blocky IMC in the filler must have created plastic strain localization at the IMC/matrix interface.