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Article

Microstructure and Properties of TA2 Titanium Joints Brazed with Ti–Zr–Cu–Ni Filler Metal

by
Zimeng Xiao
1,
Huiling Zhou
1,*,
Sheng Lu
1,*,
Zexin Wang
1 and
Oleksandr Dobuvyy
2
1
School of Material Science and Engineering, Jiangsu University of Science and Technology, Zhenjiang 212100, China
2
Department of Material Science and Technology of Metals, Admiral Makarov National University of Shipbuilding Institute, 54025 Nikolaev, Ukraine
*
Authors to whom correspondence should be addressed.
Metals 2025, 15(11), 1218; https://doi.org/10.3390/met15111218
Submission received: 3 October 2025 / Revised: 24 October 2025 / Accepted: 31 October 2025 / Published: 2 November 2025
Browse Figures
Versions Notes

Abstract

TA2 titanium was brazed with a Ti–37.5Zr–15Cu–10Ni filler metal at 860–890 °C for 20 min to investigate the influence of temperature on joint properties. Raising the brazing temperature reduced residual filler in the seam center and transformed the microstructure from heterogeneous phases to a uniform α-(Ti,Zr) solid-solution matrix, accompanied by significant widening of the diffusion layer. At brazing temperatures of 890 °C, the hardness decreased to below 300 HV0.5 and became more uniform as brittle phases were suppressed. The shear strength reached a maximum of 302 MPa, and the fracture morphology exhibited characteristics of ductile fracture. Micro-electrochemical testing indicated that the joint brazed exhibited an almost uniform current distribution and significantly reduced localized corrosion. Although a small fraction of the Widmanstätten structure was observed at this temperature, it did not impair the overall mechanical performance. These findings demonstrate that a moderate increase in brazing temperature promotes elemental diffusion, alleviates brittle phase enrichment, and markedly enhances the mechanical properties and corrosion resistance of TA2 joints.

1. Introduction

Titanium alloys are widely used in critical equipment for the aerospace, chemical, and energy industries due to their low density, high specific strength, and excellent corrosion resistance [1,2,3,4]. Typical applications include heat exchangers and honeycomb panels. Conventional metallic materials, however, are prone to pitting [5] and stress corrosion [6] cracking in environments such as seawater and chloride-containing solutions, which significantly shortens their service life. In contrast, industrially pure titanium (TA2) exhibits stable performance under these conditions [7,8]. As a result, TA2 has been extensively employed in the heat transfer units of heat exchangers, where it improves thermal efficiency and ensures long-term service reliability [9,10].
Despite the remarkable advantages of titanium and its alloys in practical applications, their joining and processing still present considerable challenges. Titanium exhibits extremely high chemical reactivity at elevated temperatures, readily reacting with oxygen, nitrogen, and hydrogen to form a brittle oxygen-enriched layer (α-case) or other fragile compounds [11,12,13,14]. These products are typically hard and brittle, and therefore, they not only have weakened interfacial bonding strength but also act as crack initiation sites during subsequent service, thereby severely reducing the toughness and fatigue life of the structure. Conventional fusion welding techniques, such as tungsten inert gas (TIG) welding [15], laser welding [16], and electron beam welding [17], often involve high heat input, rapid cooling rates, and concentrated thermal stresses when applied to titanium alloys [18,19,20]. These factors readily result in porosity, cracking, coarse columnar grains, and pronounced microstructural inhomogeneity within the joints, making it difficult to meet the stringent requirements of high-performance components [21,22,23]. More critically, during fusion welding, brittle intermetallic compounds such as Ti–Fe, Ti–Cu, and Ti–Ni are prone to form within the weld pool, and their presence significantly deteriorates the mechanical properties of the joints [24]. Consequently, although fusion welding has been widely adopted for other metallic systems, traditional welding approaches fail to achieve a proper balance among strength, toughness, and reliability for titanium and its alloys, which has greatly limited their broader application.
Brazing is a joining technique in which a filler metal with a melting point lower than that of the base material wets the substrate surface under heating and achieves metallurgical bonding through interfacial diffusion and reaction [25]. Compared with fusion welding, brazing involves lower heating temperatures, which significantly reduce the extent of the heat-affected zone, suppress the formation of coarse columnar grains and severe phase transformations, and thereby lower the risk of defects such as cracks and porosity. At the same time, brazing enables the formation of a stable reaction layer at the interface, ensuring strong metallurgical bonding [26]. For titanium and its alloys, fillers containing active elements such as Ti and Zr can effectively improve wettability, promote interfacial reactions, and ensure reliable joints [19,27,28].
In recent years, Ti-based filler metals, which combine active elements with favorable wettability, have attracted considerable attention and application in the brazing of titanium alloys. Liu et al. [29] employed a Ti–37.5Zr–10Ni–15Cu filler to vacuum braze laminated TC4 structures, and found that at 905 °C with a 10 min holding time, the typical interfacial microstructure consisted of Ti-rich α-Ti, α-Ti, (Ti,Zr)2(Cu,Ni) intermetallics, and eutectoid products. Increasing the brazing temperature and holding time was shown to alleviate the detrimental effects of (Ti,Zr)2(Cu,Ni) on joint properties. Anna et al. [30] brazed TC4 using a Ti–Cu–Zr–Pd filler at approximately 880 °C, and microstructural analysis revealed that in both approaches, the seam region was homogeneous without intermetallic segregation or Kirkendall voids. Wang et al. [23] investigated the influence of brazing temperature and holding time on the interfacial microstructure and mechanical properties of TC4/Ti–Zr–Cu–Ni/TC4 joints, and reported that raising the temperature above 900 °C and extending the holding time promoted the formation of α-Ti + (Ti,Zr)2(Cu,Ni) eutectoid phases and acicular Widmanstätten structures.
Ti-based fillers can produce joints with high strength and reliable metallurgical bonding in titanium alloys [24,30]. Research to date has concentrated mainly on process optimization at elevated temperatures or on the characterization of individual joint properties [31]. Systematic investigations into interfacial microstructural evolution at relatively lower brazing temperatures remain scarce. It helps suppress excessive grain growth of the α-Ti matrix and minimizes thermal distortion, which is highly beneficial for the joining of thin-walled and precision components in heat exchangers. In contrast, brazing conducted above 900 °C often triggers a pronounced β → α phase transformation during cooling, resulting in the formation of coarse Widmanstätten structures and continuous intermetallic compound networks, which consequently reduce the toughness of the joints. A focused study on the low-temperature brazing behavior of titanium alloys with Ti–Zr–Cu–Ni fillers is therefore essential to clarify the mechanisms that govern joint microstructure and property evolution.
In this study, industrially pure titanium (TA2) was used as the base material, and Ti–37.5Zr–10Ni–15Cu filler metal was employed for vacuum brazing at relatively low temperatures below 900 °C. By systematically comparing the interfacial microstructure, reaction layer characteristics, mechanical performance, and corrosion resistance of the joints at different brazing temperatures, the regulating effect of temperature on interfacial reactions and joint properties was comprehensively elucidated.

2. Experimental

2.1. Brazing and Base Material Treatment

In this work, industrial pure titanium (TA2) was employed as the base material and machined into two sizes, 20 mm × 20 mm × 3 mm and 10 mm × 10 mm × 3 mm, to form lap-joint configurations. The filler metal was Ti–37.5Zr–15Cu–10Ni (Type 1510, MBF-5004) alloy powder (Zhengzhou Research Institute of Mechanical Engineering Co., Ltd., Zhengzhou, China), possessing good wettability and relatively low melting characteristics, with liquidus around 840 °C [32]. Brazing was carried out at 860 °C, 870 °C, 880 °C and 890 °C, with a holding time of 20 min. During the preheating stage at around 500 °C, surface contaminants and absorbed gases were removed. When the temperature approached 800 °C, solid-state diffusion between the TA2 substrate and the Ti–37.5Zr–15Cu–10Ni filler began, serving as a preparatory stage prior to melting. At approximately 840 °C, the filler alloy melted and rapidly wetted the titanium surface, filling the joint gap through capillary action. All experiments were conducted in a vacuum brazing furnace (Zhongshan Kaixuan Vacuum Technology & Engineering Co., Ltd., Zhongshan, Guangdong, China) under a vacuum level of 5 × 10−3 Pa to minimize high-temperature oxidation and contamination. The processing parameters are illustrated in Figure 1, where the heating rate was approximately 10 °C/min, with holding stages at 500 °C and 800 °C to relieve thermal stresses and ensure sufficient diffusion between the base metal and the filler alloy.
Before brazing, the surfaces of the base material to be joined were sequentially ground with 400#, 800#, 1200#, and 1500# SiC papers to remove surface oxide films and machining traces. The specimens were then ultrasonically cleaned in anhydrous ethanol for approximately 100 s, followed by cold-air drying. Approximately 0.05 g of filler powder was weighed and uniformly spread between the upper and lower base materials as the interlayer. During assembly, ceramic plates were placed on both the upper and lower sides of the specimens to ensure uniform loading and stable positioning during heating. The assembled samples were brazed in a vacuum furnace. After brazing, the samples were furnace-cooled from the holding temperature down to approximately 100 °C at a rate of about 10 °C/min, followed by natural cooling to room temperature inside the furnace. This gradual cooling process was designed to minimize thermal stress and to allow sufficient time for the eutectoid transformation of the residual melt.

2.2. Microstructural Characterization

After wire electrical discharge machining, the brazed joints were mounted, ground, and polished on the cross-section of the seam. Grinding was carried out sequentially with 400#, 800#, 1200#, 1500#, and 2000# SiC papers, followed by fine polishing using 1.5 μm and 0.5 μm diamond suspensions to obtain a mirror-like surface. The specimens were then etched with a solution of HF:HNO3:H2O = 1:2:15 (volume ratio). The interfacial microstructure and phase distribution were examined using an optical microscope (OM, Leica DM2700M, Leica Microsystems, Wetzlar, Germany) and a scanning electron microscope (SEM, FEI Quanta 650, Thermo Fisher Scientific, Hillsboro, OR, USA). Elemental distribution and compositional analysis were performed with an energy-dispersive spectrometer (EDS, Oxford X-Max 80, Oxford Instruments, Abingdon, UK).
X-ray diffraction (XRD) was conducted using a D8 Advance diffractometer (Bruker AXS GmbH, Karlsruhe, Germany) operated at an accelerating voltage of 40 kV. The shear-fractured cross-sections of TA2 titanium alloy were fixed on the sample stage, and the measurements were performed at room temperature with a 2θ scanning range of 20–100° and a scanning rate of 5°/min. The obtained diffraction patterns were analyzed using Jade 6.0 software for phase identification and characteristic peak comparison.

2.3. Mechanical Properties

The microhardness of the brazed seam was measured using a fully automated hardness tester (HXD-1000TM/LCD, Shanghai Taiming Optical Instrument Co., Ltd., Shanghai, China) under a load of 5 N with a dwell time of 15 s. Indentations were made perpendicular to the seam at an interval of 50 μm. Nanoindentation tests were performed using a nanoindenter (Nano Indenter G200, Agilent Technologies, Santa Clara, CA, USA) with a maximum load of 20 mN, and both the loading and unloading rates were set to 40 mN/min, in order to evaluate the micromechanical properties of different phase regions.
The shear strength of the joints was determined on a universal testing machine (MTS-CMT5205, MTS Systems Co., Ltd., Shenzhen, China) at a loading rate of 0.5 mm/min. The specimen dimensions and loading configuration were prepared according to the GB/T 11363-2008 standard [33]. At least three parallel specimens were tested under each condition to ensure statistical reliability of the results.

2.4. Corrosion Resistance Testing

The localized corrosion behavior of the brazed joints was investigated using the scanning vibrating electrode technique (SVET, VersaSCAN, Ametek Princeton Applied Research, Version 1.6.6465, Oak Ridge, TN, USA). This method enabled the visualization of the local corrosion potential distribution in different joints. The testing medium was a 3.5 wt.% NaCl solution, and a platinum wire probe with a diameter of 125 μm was employed. The distance between the specimen surface and the probe tip was maintained at approximately 100 μm. A scanning area of 1 mm × 1 mm with a step size of 2 μm was selected. Prior to testing, the specimen surfaces were ground, polished, cleaned, and dried.

3. Results and Discussion

3.1. Microstructural Regions

Figure 2 shows the optical micrographs of TA2 joints brazed at four different temperatures. Overall, at all four temperatures, the filler metal exhibited good wettability and spreading behavior on the substrate surface, resulting in sound metallurgical bonding between the base material and the filler. No obvious defects such as pores, cracks, or lack of bonding were observed in the brazed seams. Based on the microstructural features, the joints can be divided into three distinct regions: Region I, the diffusion zone between the base metal and filler; Region II, the interfacial layer; and Region III, the reaction layer. With increasing brazing temperature, both the width and morphology of these three regions exhibited significant changes.
At brazing temperatures of 860 °C, 870 °C, and 880 °C, the boundaries between Region I and Region II in the joints were clearly defined, and flake-like or blocky residual phases were observed in the reaction layer (Region III). This indicates that at relatively low temperatures, the filler metal was not fully melted and uniformly diffused, leaving unreacted phases within Region III. At 890 °C, however, the microstructure of Region III became markedly denser and more homogeneous, with residual phases almost completely eliminated, and the width of the transition zone increased significantly. These phenomena can be attributed to the enhanced diffusion capability of the filler elements at elevated brazing temperatures. During brazing, Cu and Ni diffused from Region III into Region I and further into the base metal, thereby increasing the Cu and Ni content in the matrix adjacent to the seam. Since the solubility of Cu and Ni in the titanium matrix is limited, eutectoid reactions readily occurred in Regions II and III during cooling, forming a lamellar structure analogous to pearlite in steels, which consists of alternating Ti-rich and Cu/Ni-rich phases. This eutectoid structure effectively improved the density and stability of Region III. In addition, the presence of Zr suppressed the formation of continuous brittle Ti–Ni or Ti–Cu intermetallic compounds [34], which contributes to the overall toughness of the joints.

3.2. Morphology and Elemental Diffusion

As shown in Figure 3, with increasing brazing temperature, the joint morphology underwent distinct changes. The fishbone-like structures in the reaction zone gradually diminished, while the diffusion zone exhibited root-like features extending into the base metal. At brazing temperatures of 860 °C, the seam was relatively narrow, and line-scan results revealed a transition zone of approximately 152 μm. The interfacial boundaries were distinct, with Ti primarily concentrated in the base metal, while Zr, Cu, and Ni were markedly enriched in the seam center but with limited diffusion, leading to a pronounced compositional gradient across the joint. At brazing temperatures of 870 °C, the seam width increased, the transition zone expanded to 212 μm, and the interfacial reaction layer thickened, thereby improving the uniformity of the joint microstructure. At brazing temperatures of 880 °C, the seam became further densified and the interface transition appeared more diffuse; the Ti content in the seam center decreased, while the diffusion range of Cu and Ni increased significantly, resulting in a reduced compositional gradient. The transition zone at this temperature broadened to 241 μm. At brazing temperatures of 890 °C, the distribution of the filler elements in the seam region became homogeneous with nearly no abrupt compositional changes, indicating the most sufficient interdiffusion, and the width of the transition zone reached approximately 262 μm. From the perspective of diffusion kinetics, under a constant holding time of 20 min, the thickness of the diffusion layer can be approximately described as follows [35]:
δ ∝ D · t
where δ is the thickness of the diffusion layer, D is the diffusion coefficient, and t is the diffusion time. The diffusion coefficient D follows the Arrhenius relation [36], D = D0exp(−Q/RT). Therefore, increasing the temperature exponentially enhances the effective diffusion coefficient, leading to an approximate D -type growth of the thickness of Region I. Consistent with the line-scan results, raising the brazing temperature broadened and smoothed the elemental diffusion distribution within the seam. This phenomenon is not only a consequence of the combined thermodynamic and kinetic effects but also reflects the significantly enhanced penetration ability of the liquid filler metal along fast diffusion paths such as grain boundaries at elevated temperatures [37,38,39].
Such diffusion-driven homogenization not only reflects the enhanced atomic mobility at elevated temperatures but also governs the subsequent evolution of interfacial phases and, consequently, the mechanical and electrochemical behavior of the joints. The improvements in joint strength and corrosion resistance are primarily attributed to kinetic factors rather than thermodynamic equilibrium. While thermodynamics dictates the possible phase constituents in the Ti–Zr–Cu–Ni system, the actual interfacial morphology and layer thickness are strongly influenced by the diffusion and reaction rates of Cu and Ni during brazing. Within the 860–890 °C range, sufficient atomic mobility allows the formation of uniform eutectoid lamellae without excessive intermetallic coarsening. This controlled kinetic condition results in a well-balanced interface with good metallurgical bonding and minimal residual stress, leading to enhanced mechanical and electrochemical performance.
In Figure 4a–h, the SEM micrographs of local seam regions at different brazing temperatures, together with the corresponding EDS point analyses, illustrate the evolution of phase constitution and elemental distribution within the joints, as summarized in Table 1. The point analysis results indicate that Ti and Zr remained the dominant constituents, confirming that the seam region was still primarily composed of a Ti-based solid solution. In contrast, Cu and Ni, as the key filler elements, exhibited significant variations at different positions. Based on the Ti–Cu and Ti–Ni phase diagrams, three typical phases can be identified: (1) α-(Ti,Zr) solid solution, relatively deficient in Cu/Ni; (2) eutectoid structures, consisting of alternating lamellae of α-(Ti,Zr) and (Ti,Zr)2(Cu,Ni), manifested as fine light–dark striations; and (3) (Ti,Zr)2(Cu,Ni)-type intermetallic compounds, mainly corresponding to the incompletely reacted residual filler.
As shown in Figure 5, at a brazing temperature of 860 °C, elemental diffusion in the seam was limited, and the microstructure was dominated by α-(Ti,Zr) solid solution with only isolated Cu/Ni-rich regions, indicating an incomplete eutectoid reaction with residual filler remaining. At a brazing temperature of 870 °C, the diffusion of Cu and Ni was markedly enhanced, leading to the formation of fine striated or feather-like structures and activation of the eutectoid reaction. At a brazing temperature of 880 °C, periodic fluctuations in Ti/Zr and Cu/Ni contents corresponded to alternating α-(Ti,Zr) and (Ti,Zr)2(Cu,Ni), suggesting that the eutectoid reaction was nearly complete and the seam became denser and more uniform, albeit with minor untransformed areas. At a brazing temperature of 890 °C, elemental interdiffusion was most extensive, producing a homogeneous distribution of α phases and intermetallic compounds. Residual phases were largely eliminated, and the lamellar eutectoid structure was highly consistent across the seam.
The roles of Cu and Ni in titanium alloys are of critical importance. Both are β-stabilizing elements, which significantly broaden the stability domain of the β phase at elevated temperatures and therefore tend to enrich in the seam and interfacial regions during brazing [40]. During brazing, the molten Ti–Zr–Cu–Ni filler forms a Cu- and Ni-enriched β-Ti(Cu,Ni) solid solution in the seam center. With increasing brazing temperature or holding time, Cu and Ni atoms diffuse from the liquid filler toward the Ti substrate, establishing gradual compositional gradients across the interfacial region [41,42]. Upon cooling, the β-Ti(Cu,Ni) solid solution undergoes eutectoid decomposition, producing alternating lamellae of α-(Ti,Zr) and intermetallic compounds such as TiNi and TiCu. The diffusion extent of Cu and Ni strongly affects the morphology and thickness of this lamellar structure. A moderate diffusion rate leads to fine, evenly distributed eutectoid structures, which improve joint integrity. However, excessive Cu and Ni enrichment at the joint center promotes the continuous precipitation of brittle intermetallic networks, thereby reducing ductility and shear strength. These observations indicate that controlling the diffusion of β-stabilizing elements is critical for optimizing the microstructure and mechanical performance of Ti–Zr–Cu–Ni brazed joints [43]. Meanwhile, Zr atoms could be effectively dissolved in the Ti matrix and partially substitute Ti within intermetallic compounds, forming a solid solution. This substitution effect of Zr helps suppress the continuous precipitation of brittle phases and thereby improves the toughness of the eutectoid microstructure [32].
The diffusion mechanism is illustrated in Figure 6, which clearly depicts the stepwise evolution of the microstructure between the TA2 base metal and Ti–Zr–Cu–Ni filler during heating, diffusion, and cooling. In the initial brazing stage, the two sides of the base metal consisted mainly of equiaxed α-Ti grains, while the central filler layer remained uniform, with a distinct boundary separating them and no metallurgical interaction having yet occurred. As the temperature increased into the brazing range, Cu and Ni, as typical β-stabilizing elements, diffused from the filler side toward the base metal, whereas Ti diffused in the opposite direction into the seam, accompanied by gradual enrichment of Zr at the interface. Such bidirectional diffusion produced a pronounced compositional gradient near the interface, providing the conditions for subsequent reaction phases to form. During the holding stage, solid-state reactions were initiated at the interface, generating lamellar structures composed of alternating α-(Ti,Zr) solid solution and (Ti,Zr)2(Cu,Ni) intermetallic compounds. These eutectoid phases appeared as fine light–dark striations in SEM images, while residual unreacted filler could still be observed in some regions. With further heating and prolonged diffusion, these alternating lamellae gradually became continuous, and the seam evolved from an initially heterogeneous reaction zone into a dense and stable eutectoid microstructure. Upon cooling, the β solid solution fully decomposed into alternating lamellae of α and intermetallic compounds, while occasional Widmanstätten α appeared, indicating complex transformation kinetics. With sufficient diffusion at elevated temperatures, the seam evolved into a dense and uniform lamellar eutectoid structure characterized by blurred interfaces and smooth elemental transitions, underscoring the synergistic role of elemental diffusion and phase transformation [44].

3.3. Mechanical Performance of the Brazed Joints

The brazing temperature exerted a significant influence on the microhardness of the seam region. As shown in Figure 6, the hardness distribution of the joints exhibited a distinct trend with increasing brazing temperature. At 860 °C, the reaction layer at the seam center exhibited the highest hardness, reaching 535 HV0.5. This was attributed to the enrichment and continuous formation of intermetallic compounds such as TiCu and Ti2Ni, which possess high bonding energy and brittle crystal structures, thereby markedly increasing the local hardness. In addition, both the interfacial layer and the diffusion zone also showed elevated hardness values, which can be closely related to the presence of fine lamellar eutectoid structures, where lamellar interfaces and grain refinement provided additional strengthening. With further increases in brazing temperature, the hardness of the reaction and interfacial layers decreased significantly. This reduction can be explained by enhanced diffusion of Cu and Ni into the Ti matrix under higher thermal input, which weakened the continuity of intermetallic compounds and reduced the volume fraction of brittle eutectoid phases. Hardness contour maps further demonstrated that at 890 °C, the overall hardness distribution across the joint became more uniform, mainly due to the transformation of the seam microstructure into α-(Ti,Zr) solid solution structure. Therefore, appropriately raising the brazing temperature can effectively suppress the formation and continuity of hard and brittle phases at the interface.
Nanoindentation tests were performed at the center of the brazed seam, as shown in Figure 7. With increasing brazing temperature, the nanoindentation values obtained from three randomly selected points at the seam center gradually converged, indicating that the microstructure evolved from a multiphase coexistence to a more homogeneous structure dominated by a single solid solution. At 860 °C, the seam center contained a large fraction of continuous and dense (Ti,Zr)2(Cu,Ni) intermetallic compounds together with fine lamellar eutectoid structures, resulting in a pronounced increase in local hardness and strong heterogeneity. When the temperature rose to 870 °C, the volume fraction of intermetallic compounds decreased significantly, and their continuity was disrupted. Most testing points were located in eutectoid regions, leading to an overall reduction in hardness and a smaller degree of fluctuation. With further temperature increase, the seam center became progressively dominated by an α-(Ti,Zr) solid solution, while the amounts of intermetallic compounds and eutectoid phases were greatly reduced. As a result, the hardness values decreased to below 6 GPa and tended to become uniform. Therefore, the rise in brazing temperature not only reduced the amount and continuity of hard and brittle phases in the seam but also promoted the formation and stabilization of the α-(Ti,Zr) solid solution. This evolution led to a more balanced hardness distribution at the local scale and reflected the transition of mechanical behavior from brittleness to toughness, which is of great significance for improving the overall reliability of the joints.
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Figure 6. Microhardness distributions of TA2 joints brazed at different temperatures: (a) comparison of hardness in the reaction layer, interface layer, and diffusion layer; (b) hardness contour maps of joint cross-sections.
Figure 6. Microhardness distributions of TA2 joints brazed at different temperatures: (a) comparison of hardness in the reaction layer, interface layer, and diffusion layer; (b) hardness contour maps of joint cross-sections.
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Figure 7. Nanoindentation load–displacement curves and corresponding hardness values of different phases in the seam center of TA2 joints brazed at various temperatures: (a) 860 °C; (b) 870 °C; (c) 880 °C; (d) 890 °C.
Figure 7. Nanoindentation load–displacement curves and corresponding hardness values of different phases in the seam center of TA2 joints brazed at various temperatures: (a) 860 °C; (b) 870 °C; (c) 880 °C; (d) 890 °C.
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Figure 8 shows the effect of brazing temperature on joint shear strength. The strength increased from 186 MPa at a brazing temperature of 860 °C to 302 MPa at 890 °C. The load–displacement curves indicated that the joint brazed at 890 °C exhibited the highest peak load and fracture displacement, while the 860 °C joint fractured prematurely with inferior strength and ductility. Macroscopic fracture morphologies further confirmed this trend: low-temperature joints showed small cleavage-dominated fracture areas with cracks induced by residual filler and continuous intermetallic compounds, whereas high-temperature joints displayed enlarged ductile fracture regions with suppressed crack propagation. These findings are consistent with the observed microstructural evolution, where elevated temperature promoted Cu/Ni diffusion and homogenization of the reaction layer, thereby reducing brittle phases and enhancing interfacial bonding, ultimately leading to optimal shear strength.
The evolution of fracture morphology further revealed the influence of brazing temperature on the fracture mechanism of the joints. As shown in Figure 9, with the brazing temperature increasing from 860 °C to 890 °C, the fracture characteristics gradually transitioned from typical brittle cleavage to ductile fracture dominated. At lower temperatures, the fracture surface was mainly composed of large cleavage facets and river patterns, accompanied by through-thickness cracks, indicating stress concentration and brittle failure caused by the continuous enrichment of intermetallic compounds. When the temperature was raised to 870 °C, the fracture surface was still dominated by quasi-cleavage features, but shallow dimples and tear ridges began to appear locally. This suggests that the reduced continuity of brittle phases deflected and blunted the crack propagation path, resulting in a mixed fracture mode with both brittle and ductile features. At 890 °C, the fracture was primarily composed of uniform and deep equiaxed dimples, with cleavage traces almost absent. This evolution is closely related to the microstructural features: at lower temperatures, residual filler metal and Cu/Ni-rich regions tended to form continuous brittle phases, serving as weak points for rapid crack initiation and propagation; whereas at higher temperatures, the homogenized diffusion of elements and the stable formation of the α-(Ti,Zr) solid solution effectively suppressed the continuity of intermetallic compounds, reduced potential crack sources, and promoted crack growth through energy-consuming ductile mechanisms, thereby enabling the transition from brittle to ductile fracture. The XRD results of the fracture surfaces are presented in Figure 10. In all samples, α-Ti matrix phases along with a certain amount of CuTi2, CuZr2, and NiTi were detected. With increasing brazing temperature, both the number and intensity of the diffraction peaks decreased progressively.

3.4. Corrosion Properties

As shown in Figure 11, the micro-electrochemical test results of the joints clearly illustrate the effect of brazing temperature on local corrosion behavior. At 860 °C, the current distribution exhibited large alternating regions of high positive and negative values. This indicates the presence of pronounced micro-galvanic effects within the seam, where corrosion was mainly concentrated at the phase boundaries between α-(Ti,Zr) and Cu/Ni-rich intermetallic compounds [45]. The segregation induced by residual filler solidification further aggravated the non-uniformity of corrosion. In contrast, the joint brazed at 890 °C displayed a nearly uniform current distribution with only a few scattered weak hotspots, and the overall corrosion activity was significantly reduced, indicating optimal electrochemical stability.
This difference can be attributed to diffusion and reaction mechanisms. At lower brazing temperatures, limited Cu and Ni diffusion resulted in residual filler at the seam center, which reacted with the Ti matrix to form continuous intermetallic compounds. The potential difference between these compounds and the α-(Ti,Zr) matrix induced strong current localization and severe micro-galvanic corrosion. At higher brazing temperatures, enhanced Cu and Ni diffusion minimized residual filler and reduced intermetallic compound formation, thereby diminishing both the potential difference and interfacial density. In addition, Zr in solid solution promoted the formation of a stable and compact TiO2–ZrO2 passive film [46], further lowering local corrosion susceptibility. Thus, increasing the brazing temperature not only suppressed the continuity of brittle phases but also reinforced the protective role of the passive film, leading to the disappearance of micro-current field heterogeneity and a marked improvement in joint corrosion resistance.

4. Conclusions

In this study, stable brazing of TA2 titanium was achieved at relatively low temperatures using a Ti–37.5Zr–15Cu–10Ni filler metal. The joints exhibited dense and continuous interfaces without macroscopic defects such as pores or cracks. The microstructure and properties of the brazed joints were systematically analyzed, and the following conclusions were shown:
  • With increasing brazing temperature, the seam microstructure evolved from residual filler and heterogeneous phases to a uniform structure characterized by smooth elemental distribution and gradual interfacial transition. The seam center was mainly composed of α-(Ti,Zr) solid solution, (Ti,Zr)2(Cu,Ni) intermetallic compounds, and lamellar eutectoid structures. At 890 °C, elemental diffusion was sufficient, and the microstructure transformed into an α-(Ti,Zr)-dominated solid solution with significantly improved homogeneity and stability.
  • Mechanical tests revealed that low-temperature joints exhibited higher and non-uniform hardness due to the enrichment of hard brittle phases, and the shear strength was only 186 MPa. As the brazing temperature increased, the overall hardness decreased and became more uniform, while the shear strength increased progressively to 302 MPa at 890 °C. The fracture mechanism transitioned from cleavage-dominated brittle fracture to ductile fracture with equiaxed dimples.
  • Low-temperature joints exhibited strong anodic hotspots and pronounced micro-galvanic corrosion, mainly concentrated at α-(Ti,Zr)/IMCs phase boundaries. With increasing brazing temperature, both the number of hotspots and current density decreased, and at 890 °C, the current distribution became nearly uniform, demonstrating markedly improved corrosion resistance. The underlying mechanism was attributed to sufficient elemental diffusion and the formation of a protective TiO2–ZrO2 composite film at higher temperatures, which jointly mitigated electrochemical heterogeneity and enhanced interfacial stability.

Author Contributions

Conceptualization, H.Z., O.D. and S.L.; Methodology, Z.X.; Software, Z.W.; Validation, Z.X.; Investigation, H.Z.; Resources, O.D. and S.L.; Data Curation, O.D.; Writing—Original Draft Preparation, Z.X.; Writing—Review and Editing, H.Z.; Supervision, Z.W.; Funding Acquisition, H.Z. and Z.W. All authors have read and agreed to the published version of the manuscript.

Funding

This work is supported by the National Key Research and Development Program of China (Grant No. 2021YFB3401100) and the Jiangsu Province Foreign Experts Hundred Talents Program (BX2022030).

Data Availability Statement

The original contributions presented in this study are included in the article. Further inquiries can be directed to the corresponding authors.

Conflicts of Interest

The authors declare no conflicts of interest.

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Figure 1. Schematic illustration of the brazing experiment for TA2 titanium alloy: (a) assembly configuration of the specimens; (b) process parameters for vacuum brazing.
Figure 1. Schematic illustration of the brazing experiment for TA2 titanium alloy: (a) assembly configuration of the specimens; (b) process parameters for vacuum brazing.
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Figure 2. Optical micrographs of TA2 joints brazed at different temperatures: (a) 860 °C; (b) 870 °C; (c) 880 °C; (d) 890 °C.
Figure 2. Optical micrographs of TA2 joints brazed at different temperatures: (a) 860 °C; (b) 870 °C; (c) 880 °C; (d) 890 °C.
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Figure 3. Interfacial microstructures and elemental distributions of TA2 joints brazed at (a) 860 °C (b) 870 °C (c) 880 °C and (d) 890 °C: (a1–d1) SEM images of joint cross-sections; (a2–d2) elemental line scans of Ti, Zr, Cu, and Ni; (a3–d3) element distribution map.
Figure 3. Interfacial microstructures and elemental distributions of TA2 joints brazed at (a) 860 °C (b) 870 °C (c) 880 °C and (d) 890 °C: (a1–d1) SEM images of joint cross-sections; (a2–d2) elemental line scans of Ti, Zr, Cu, and Ni; (a3–d3) element distribution map.
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Figure 4. Local magnified morphologies of TA2 joints brazed at different temperatures: (a,b) 860 °C; (c,d) 870 °C; (e,f) 880 °C; (g,h) 890 °C.
Figure 4. Local magnified morphologies of TA2 joints brazed at different temperatures: (a,b) 860 °C; (c,d) 870 °C; (e,f) 880 °C; (g,h) 890 °C.
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Figure 5. Schematic illustration of interfacial evolution during brazing: (a) initial stage; (b) elemental diffusion stage; (c) reaction phase formation; (d) joint formation.
Figure 5. Schematic illustration of interfacial evolution during brazing: (a) initial stage; (b) elemental diffusion stage; (c) reaction phase formation; (d) joint formation.
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Figure 8. Shear performance of TA2 joints brazed at different temperatures: (a) load–displacement curves; (b) macroscopic fracture morphologies after shear testing.
Figure 8. Shear performance of TA2 joints brazed at different temperatures: (a) load–displacement curves; (b) macroscopic fracture morphologies after shear testing.
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Figure 9. Fracture morphologies of TA2 joints brazed at different temperatures: (a,b) 860 °C; (c,d) 870 °C; (e,f) 880 °C; (g,h) 890 °C.
Figure 9. Fracture morphologies of TA2 joints brazed at different temperatures: (a,b) 860 °C; (c,d) 870 °C; (e,f) 880 °C; (g,h) 890 °C.
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Figure 10. XRD patterns of the fracture surface of the brazed joints.
Figure 10. XRD patterns of the fracture surface of the brazed joints.
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Figure 11. SVET current distribution maps of the seam region in TA2 joints brazed at different temperatures: (a) 860 °C; (b) 870 °C; (c) 880 °C; (d) 890 °C.
Figure 11. SVET current distribution maps of the seam region in TA2 joints brazed at different temperatures: (a) 860 °C; (b) 870 °C; (c) 880 °C; (d) 890 °C.
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Table 1. EDS point analysis results of the marked positions in Figure 5.
Table 1. EDS point analysis results of the marked positions in Figure 5.
PointsTiZrCuNiPossible Phases
172.1814.779.763.29α-(Ti,Zr)
282.271.952.0213.76eutectoid phases
368.6021.356.883.18α-(Ti,Zr)
432.3241.1015.1511.43(Ti,Zr)2(Cu,Ni)
568.7119.097.354.85α-(Ti,Zr)
665.8822.576.155.41α-(Ti,Zr)
770.0710.0211.198.72eutectoid phases
831.4241.5116.2910.77(Ti,Zr)2(Cu,Ni)
930.9542.6314.8411.59(Ti,Zr)2(Cu,Ni)
1067.6319.516.756.10α-(Ti,Zr)
1176.6617.093.482.77α-(Ti,Zr)
1266.149.7414.619.51eutectoid phases
1334.8239.2915.4410.45(Ti,Zr)2(Cu,Ni)
1476.0812.408.872.65α-(Ti,Zr)
1568.228.9113.459.41eutectoid phases
1660.9119.7615.054.28eutectoid phases
1737.8735.5411.6314.96(Ti,Zr)2(Cu,Ni)
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Xiao, Z.; Zhou, H.; Lu, S.; Wang, Z.; Dobuvyy, O. Microstructure and Properties of TA2 Titanium Joints Brazed with Ti–Zr–Cu–Ni Filler Metal. Metals 2025, 15, 1218. https://doi.org/10.3390/met15111218

AMA Style

Xiao Z, Zhou H, Lu S, Wang Z, Dobuvyy O. Microstructure and Properties of TA2 Titanium Joints Brazed with Ti–Zr–Cu–Ni Filler Metal. Metals. 2025; 15(11):1218. https://doi.org/10.3390/met15111218

Chicago/Turabian Style

Xiao, Zimeng, Huiling Zhou, Sheng Lu, Zexin Wang, and Oleksandr Dobuvyy. 2025. "Microstructure and Properties of TA2 Titanium Joints Brazed with Ti–Zr–Cu–Ni Filler Metal" Metals 15, no. 11: 1218. https://doi.org/10.3390/met15111218

APA Style

Xiao, Z., Zhou, H., Lu, S., Wang, Z., & Dobuvyy, O. (2025). Microstructure and Properties of TA2 Titanium Joints Brazed with Ti–Zr–Cu–Ni Filler Metal. Metals, 15(11), 1218. https://doi.org/10.3390/met15111218

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