Effects of Ti Addition on Microstructure, Mechanical Properties and Corrosion Resistance of the Cu-Zn-Ni Alloy

Abstract

The effects of Ti addition on the microstructure, mechanical properties and corrosion resistance of the Cu-30Zn-6Ni-1.5Ti (wt.%) alloy were investigated in the present study. Microstructure analysis confirms that the Ni₃Ti and NiTi phases are formed in the Cu-Zn-Ni-Ti alloy. Most of the Ni₃Ti particles dissolve into the matrix, whereas NiTi particles remained after the solution treatment. Moreover, nano-sized Ni₃Ti and Cu₂NiZn phases are precipitated from the matrix during the aging process. The yield strength improvement of the studied alloy is attributed to the Orowan strengthening (accounting for 42.5% of the total yield strength), then the grain boundary strengthening (31.1%), the dislocation strengthening (16.9%), the twin boundary strengthening (5.0%) and the solid solution strengthening (0.5%). After cold rolling with 90% deformation and aging at 440 °C for 1 h, the designed alloy has a hardness of 327.1 HV, a yield strength of 1192 MPa, an ultimate tensile strength of 1297 MPa, an elongation of 1.8%, an electrical conductivity of 11.2% IACS and a corrosion rate of 2.8 μm/month in 3.5% NaCl solution.

1. Introduction

Copper-based alloys have been widely used in the fields of electronic connectors [1,2,3], heat conductors [4,5], heat exchangers [6] and seawater applications [7,8] because of their excellent combination of mechanical properties, electrical conductivity, thermal stability and corrosion resistance. Cu-Zn-Ni alloys, known as “nickel-brass” [9], exhibit excellent comprehensive properties among mechanical properties, electrical conductivity and corrosion resistance, making them a competitive candidate material for the above applications. Zn plays an important role in decreasing the melting point and minimizing the porosity, whereas Ni is responsible for enhancing the mechanical properties and improving the corrosion resistance [10]. According to the ternary Cu-Zn-Ni diagrams reported in the studies [11,12], Zn and Ni generally have a large solid solubility in copper, thus the improvement of mechanical properties is limited and the electrical conductivity is significantly damaged.

In order to solve such difficulties, numerous attempts have been made by microalloying [13,14,15]. Especially, precipitation strengthening by adding a fourth component can realize the simultaneous improvement of mechanical properties and the electrical conductivity, and it is a more common and feasible way to enhance the comprehensive performance of Cu-Zn-Ni alloys. Zhou et al. [13,14] discovered that the mechanical properties and electrical conductivity of the Cu-20Zn-10Ni-1.2Al (wt.%) alloy were dramatically improved by aging at 500 °C for 1 h due to the formation of nano-sized precipitates with the L1₂ type ordered structure. Chen et al. [15] developed a Cu-10Zn-1.5Ni-0.34Si (wt.%) alloy and obtained a UTS of 650 MPa and an electrical conductivity of 28.5% IACS due to the introduction of δ-Ni₂Si precipitates. However, the above studies mainly pay attention to the improvement of mechanical properties and electrical conductivity, and few studies have been focused on the corrosion behavior of the alloys. Considering the harsh work environment of Cu-Zn-Ni alloys in the practical applications, the corrosion resistance is also an important factor to be taken into consideration [16].

Ti-based alloys are important structural materials used for industrial applications due to their excellent properties such as high mechanical properties [17], superior strength-to-weight ratio [18] and good corrosion resistance, especially in corrosive environments containing chloride ions [17,19]. For this reason, an attempt was made to add Ti, as a trace element, to Cu-Ni-based alloys, and exciting results have been obtained [20,21]. Ti has a strong affinity to Ni atoms to form Ni_xTi intermetallic compounds, which decreases the solid solubility of Ni in the Cu matrix and thereby causes an improvement in the tensile strength and electrical conductivity. However, little research on the effect of Ti addition on the microstructure, precipitation behavior and properties of the Cu-Zn-Ni alloy has been reported.

Considering the aforementioned positive effects of Ti addition in the Cu-Ni-based alloys, it is reasonable to design a new type of Cu-Zn-Ni-Ti alloy in the present study. The main aim of this work is to investigate the effects of Ti addition on the properties of the Cu-Zn-Ni alloy in terms of microstructure, mechanical properties and corrosion resistance.

2. Materials and Methods

2.1. Material Preparation

In this study, two alloys were prepared in an 8 kg intermediate frequency induction furnace using electrolytic copper (99.97 wt.%), zinc (99.8 wt.%), nickel (99.9 wt.%), and titanium (99.9 wt.%) as raw materials. The Cu and Ni were melted in the furnace, and Ti was added to the melt when the temperature reached 1400 °C. Effective dissolution of high-melting-point Ti particles in the low-melting-point Cu-Ni alloy can be achieved through appropriate temperature control and alloy composition design. After the complete melting of Ti, Zn was added to the melt at a temperature of 1000 °C. The melt was then poured into a graphite mold of 300 × 30 × 200 mm³ dimension and a preheated temperature of 600 °C. The designed compositions of the alloys and their actual compositions tested by the X-ray fluorescence spectrum technique (XRF-1800) are presented in Table 1.

In order to improve their mechanical properties and break up coarse casting microstructures, the as-cast ingots were homogenized at 850 °C for 2 h, then quickly hot rolled to 18 mm. After solution treatment at 880 °C for 2 h followed by a rapid water-quenching, the samples were machined on both sides to remove surface defects and cold rolled (CR) with 90% deformation. Finally, the samples were, respectively, aged at 440 °C for various times and aged at different temperatures for 1 h.

2.2. Microstructural Analysis

The phase identification of the samples was performed using an X-ray diffractometer (XRD) equipped with a Cu target operating at 1.8 kW, and a series of θ–2θ scans were conducted to provide a record of the XRD patterns at room temperature. Dislocation densities were calculated by the Williamson–Hall method [22,23], which can be described as follows:

βcos(θ) = 4εsin(θ) + Kλ/d (1)

(2)$\rho ={\displaystyle \frac{2\sqrt{3}\epsilon }{\mathrm{db}}}$

Here, β is the full width at half maximum (FWHM) of the Cu diffraction peak (in radians), θ is the diffraction angle, K is a constant related to the instrument and sample, λ is the wavelength of the X-rays, d is the grain size, ρ is the dislocation density, ε is the micro-strain computed from the true XRD result, and b is the Burgers vector. As an example, Figure 1 shows the relation between peak broadening βcos(θ) and 4sin(θ) for the Cu-Zn-Ni-Ti alloy after CR 90% deformation and aging at 440 °C for 1 h. The value of micro-strain ε is deduced from the slope of the fitted line.

The microstructures of the samples were observed by an optical microscope (OM, Olympus GX51, Tokyo, Japan) and a field emission scanning electron microscope (SEM, Zeiss Supra 55, Oberkochen, Germany) operating at an acceleration voltage of 15 kV. For transmission electron microscopy (TEM), discs of 3 mm in diameter were first punched then ground to 30 µm and, finally, ion-thinned by ion milling (PIPS 691). TEM and high-resolution TEM (HRTEM) were performed on an FEI Tecnai G220 transmission electron microscope (Eindhoven, The Netherlands) operating at an acceleration voltage of 200 kV.

2.3. Mechanical Properties and Electrical Conductivity Tests

Vickers hardness tests were carried out on an MH5L-type microhardness tester under a load of 300 g and loading time of 15 s. Tensile strength tests were performed at room temperature on the SANS universal testing machine (UTM5150), with an initial displacement rate of 2 mm/min. Settling cycles were applied prior to the tensile tests. In addition, Each specimen underwent three repeated tensile tests, and the average value was taken to ensure the accuracy of tensile results. The dimension of the tensile specimens is shown in Figure 2. Electrical conductivity was measured by a D60K conductivity measuring instrument, according to the International Annealed Copper Standard (%IACS) [24].

2.4. Corrosion Rate Tests

Samples for the mass loss measurement were cut into a dimension of 20 × 20 × 5 mm³. Each sample was ground by SiC abrasive papers from 240 grit to 1000 grit, ultrasonically cleaned in ethanol for 5 min, rinsed with distilled water and dried in cold air. After weighing carefully on an analytical balance (ME204E, Mettler Toledo Corp., Greifensee, Switzerland) with an accuracy of ±0.1 mg, each sample was immersed into a container with 300 mL of 3.5% NaCl solution, and the exposed area of each sample was calculated independently. The temperature of the solution was maintained at 25 ± 1 °C throughout the entire experimental procedure in a thermostatic water bath. In order to compensate for the evaporation loss, an appropriate amount of distilled water was added to the solution every day to maintain the original volume. Furthermore, the solution was replaced with a fresh one every week. At the end of the corrosion test, the samples were immersed in a 1:1 HCl water solution for 3 min to remove the corrosion products on the surfaces, rinsed with distilled water, dried in cold air and put into a thermostatic drier box at 60 °C for 2 h, then finally, weighed.

Corrosion samples with a dimension of 10 × 10 × 5 mm³ were prepared for the surface morphology analysis. These samples were successively ground with 240–2000 grit SiC papers and polished with a 1.5 μm diamond paste. After immersing the samples in 3.5% NaCl solution for 30 days, the corroded surface morphologies of the samples were observed by SEM.

2.5. Electrochemical Tests

Potentiodynamic polarization (PDP) was adopted to investigate the corrosive behavior of the samples using on a electrochemical workstation (CS2350H, Wuhan Kest Instrument Co., Ltd., Wuhan, China) equipped with a typical three-electrode cell system, where the saturated calomel electrode (SCE) worked as a reference electrode, the platinum electrode acted as an auxiliary electrode, and the selected samples served as the working electrode. The electrolyte was 3.5 wt.% NaCl solution. The electrode specimens with an exposed area of 15 × 15 mm² were first ground by SiC papers from 240 to 2000 grit size, then the surface was polished and instantly dried in cold air. The PDP tests were administered in the voltage range from −0.6 to 2 V/SCE at a scan rate of 1 mV/s. The corrosion parameters of polarization curves, including the corrosion potential (E_corr), corrosion current density (J_corr) and the Tafel slopes of anodic and cathodic (${\beta }_{\mathrm{a}}$ and ${\beta }_{\mathrm{c}}$) values, were fitted and calculated using EC-LabV10.40 software.

3. Results and Discussion

3.1. Microstructure

3.1.1. Microstructure of the As-Cast Alloys

Figure 3 exhibits the optical micrographs of the as-cast Cu-Zn-Ni and Cu-Zn-Ni-Ti alloys. It can be observed from Figure 3a that a large number of uniform equiaxed grains with an average grain size of ~68 µm are distributed in the matrix. A few micrometer-sized particles with an irregular morphology are distributed within the grains and along the grain boundaries (GBs). With Ti addition (Figure 3b), a dendritic microstructure is formed, which is completely different from the Cu-Zn-Ni alloy. In addition, a sharp change in the phase morphology is also noticed, and different types of newly-formed intermetallic compounds are dispersed inside the alloy. The results above suggest that the addition of the Ti element influences the solid–liquid interface dynamics and promotes the formation of dendrites. In addition, during solidification, the interdendritic spaces are the last to solidify and trap higher concentrations of alloying elements due to solute segregation, leading to the preferential formation of intermetallic compounds within or adjacent to the dendrite arms.

In order to clarify the phase constitution of the two alloys, the secondary electron scanning and XRD analysis were performed. Figure 4 presents the SEM images of the as-cast Cu-Zn-Ni and Cu-Zn-Ni-Ti alloys, and the corresponding energy dispersive spectrometer (EDS) results are illustrated in Table 2. For the Cu-Zn-Ni alloy (Figure 4a), most Zn and Ni atoms dissolve in the matrix, and a small amount of Ni-rich particles segregate within the grains and along the grain boundary (marked by points A and B). With the addition of Ti (Figure 4b), two different phases coexist in the matrix. One kind presents an irregular morphology with an average size of ~4.1 µm, and the other exhibits a spheroidal morphology with a size of several tens of nanometers. The coarse particles with an irregular morphology are ascertained as the NiTi phase (point D). However, the crystal structure of the nano-sized phase cannot be accurately detected (point E and F) due to its small size. Thus, further studies are needed to reveal the phase structure.

Figure 5 illustrates the XRD patterns of the two alloys to identify the phases. Typical peaks of α-Cu can be identified in both cases, and the weak peaks of the NiTi phase are also detected with Ti additions (Figure 5a). However, the other phase still cannot be detected accurately due to the small volume fraction. For a more detailed examination, the magnified patterns between 34° and 52° are presented in Figure 5b, and the additional weak peaks are identified and indexed as the Ni₃Ti phase.

3.1.2. Microstructures of the Alloys After Solution Treatment

Figure 6 illustrates the microstructures of the two alloys subjected to the solution treatment. It can be observed from Figure 6a that the grains significantly grow up for both of the alloys during solution treatment process. The addition of Ti has a significant effect on grain refinement and reduces the average grain size from ~216 µm to ~172 µm compared with the Cu-Zn-Ni alloy (Figure 6b). The single phase can be clearly observed in Figure 6c, indicating that the Ni-rich phase has completely dissolved into the matrix for the Cu-Zn-Ni alloy. With Ti additions, most of the Ni₃Ti particles with a smaller size are dissolved in the matrix, whereas a small amount of Ni₃Ti particles with a larger size and NiTi particles remain (Figure 6d). These particles can act as barriers for grain boundary migration by Zener pinning [25]. According to the Zener’s equation, the pinning pressure of the particles can be expressed as [26]

(3)$Z_P={\displaystyle \frac{3{\gamma }_{DB}{F}_V}{2r}}$

Here, γ_DB is the energy per unit area for a high-angle boundary, r is the particle radius, and F_v is the particle volume fraction. For the Cu-Zn-Ni-Ti alloy, the NiTi particles mainly lie inside grains, and the partially undissolved Ni₃Ti particles segregate at grain boundaries. It can be inferred that during the solution treatment process, grain boundaries easily move across the NiTi particles and are pinned by the Ni₃Ti particles with a much smaller size. Hence, the Cu-Zn-Ni-Ti alloy reduces the mobility of grain migration and exhibits a better refining effect compared to the Cu-Zn-Ni alloy.

3.1.3. Microstructures of the Alloys After Cold Rolling

The SEM images of the Cu-Zn-Ni and Cu-Zn-Ni-Ti alloys with CR 90% deformation are displayed in Figure 7. It can be observed from Figure 7a that no secondary phase is formed in the Cu-Zn-Ni alloy, and the coarse grains bend toward the rolling direction and evolve into lamellar structures. For the Cu-Zn-Ni-Ti alloy (Figure 7b), Ni₃Ti and NiTi particles are unevenly distributed along the rolling direction. Grain boundaries between the lamellar structures are narrowed and are difficult to distinguish, which indicates a higher dislocation density compared with the Cu-Zn-Ni alloy. The undissolved Ni₃Ti and NiTi particles impede the movement of dislocations and cause the multiplication of dislocations during the CR process, thereby significantly increasing the dislocation density.

3.1.4. Microstructures of the Alloys After Aging

The precipitated phase structure, size and distribution of the alloys during the aging process determine the final mechanical properties. Thus, the TEM analysis was carried out to investigate the structure and distribution of precipitates for the two alloys.

Figure 8 illustrates the high-magnification TEM images of the Cu-Ni-Zn alloy with CR 90% deformation and aging at 440 °C for 1 h. It can be observed from Figure 8a that the high density of dislocations and a small amount of nano-sized precipitates tangle together to form dislocation cells. In order to clarify the phase structure of precipitates, HRTEM and corresponding fast Fourier transformation (FFT) images are carried out (Figure 8b). The obtained results suggest that the precipitates with an average diameter of ~5 nm are the cubic Cu₂NiZn phase, with the cell parameters of a = 0.4588 nm, b = 0.4588 nm and c = 0.4588 nm. The precipitated phase structure is coherent with the α-Cu matrix and is formed during the aging process, which is consistent with the results reported in Refs. [10,14]. Moreover, a large number of subgrains and deformation twins also exist in the alloy (marked by the red arrows in Figure 8c,d). The average grain size and twin width measured from 20 images are ~117 nm and ~36 nm, respectively.

High-magnification TEM images of the Cu-Ni-Zn-Ti alloy with CR 90% deformation and aging at 440 °C for 1 h are exhibited in Figure 9. Unlike the Cu-Zn-Ni alloy, a spherical particle is detected in the Cu-Zn-Ni-Ti alloy (Figure 9a). The corresponding SAED pattern shown in Figure 9b confirms that it is the cubic NiTi phase with the cell parameter of a = 0.4558 nm, which is in accordance with the EDS and XRD results in Table 2 and Figure 3. Figure 9c,e suggest that two different types of nano-sized precipitates (marked by red arrows) exist in the matrix. The HRTEM and corresponding FFT images of the marked area in the inset of Figure 9d suggest that the precipitates with an average diameter of ~15 nm can be ascertained as hexagonal Ni₃Ti phase, with the cell parameters of a = 0.5092 nm, b = 0.5092 nm and c = 0.8297 nm. The finding is consistent with the result reported by Liu et al. [21]. Moreover, the other precipitates with an average diameter of ~5 nm correspond to the cubic Cu₂NiZn phase (Figure 9f). It is noticeable that the number of Cu₂NiZn precipitates in the Cu-Zn-Ni-Ti alloy is much more than that of the Cu-Zn-Ni alloy after aging. This is due to the higher dislocation density introduced in the CR process (Figure 7), which provides more nucleation sites for the precipitation of the precipitates during the aging process. In addition, similar to the Cu-Zn-Ni alloy, the subgrains and deformation twins are also observed in the Cu-Zn-Ni-Ti alloy (Figure 9a,e) with an average size of ~106 nm and width of ~31 nm.

3.2. Mechanical Properties and Electrical Conductivity

3.2.1. Hardness and Electrical Conductivity

In order to determine the optimum aging parameters for the superior overall performance of the Cu-Zn-Ni and Cu-Zn-Ni-Ti alloys, the variations in hardness and electrical conductivity of the alloys with CR 90% deformation and aging at 440 °C for various times and aging at different temperatures for 1 h are investigated.

Figure 10 exhibits the variations in hardness and electrical conductivity of the alloys with CR 90% deformation and aging at 440 °C for various times. It can be seen from Figure 10a that the hardness of the Cu-Zn-Ni alloy slightly increases at the early stage of aging and then falls in a sharp decrease after aging for 1 h. In contrast, the hardness of the Cu-Zn-Ni-Ti alloy increased greatly in the initial aging stage and then slowly decreased. For the Cu-Zn-Ni-Ti alloy, the highest hardness value of 323.1 HV is obtained after aging at 440 °C for 1 h, whereas for the Cu-Zn-Ni alloy, the value is only 244.7 HV. In addition, the Cu-Zn-Ni-Ti alloy exhibits a higher softening resistance compared to the Cu-Zn-Ni alloy. Even after aging for 10 h, the hardness of the Cu-Zn-Ni-Ti alloy only decreases by 14.9% compared with the corresponding peak hardness, while that is 26.1% for the Cu-Zn-Ni alloy. The variations in electrical conductivity of the alloys with CR 90% deformation and aging at 440 °C for various times are presented in Figure 10b. Before aging, the electrical conductivity of the Cu-Zn-Ni-Ti alloy is 1.1% IACS lower than that of the Cu-Zn-Ni alloy due to the existence of partial Ti atoms in the matrix. With the increase in aging time, the electrical conductivity of the Cu-Zn-Ni increases slowly, while the electrical conductivity increases rapidly in the initial aging and approaches a stable value after aging for 1 h for the Cu-Zn-Ni-Ti alloy. The equilibrium electrical conductivity of the Cu-Zn-Ni-Ti alloy is 11.2% IACS, which is 2.4% IACS higher than that of the Cu-Zn-Ni alloy. The higher hardness and electrical conductivity of the Cu-Zn-Ni-Ti alloy are mainly attributed to the formation of Ni₃Ti and Cu₂NiZn precipitates from the supersaturated solid solution during the aging process. These precipitates significantly increase the precipitation strengthening and decrease the impurity scattering, thereby improving the mechanical properties and the electrical conductivity.

The variations in hardness and electrical conductivity of the alloys with CR 90% deformation and aging at different temperatures for 1 h are exhibited in Figure 11. It can be observed from Figure 11a that no dramatic improvement occurs in the hardness for the Cu-Zn-Ni alloy, but a rapid decrease is detected when the temperature is higher than 440 °C. The Cu-Zn-Ni-Ti alloy exhibits a higher precipitation strengthening response, which reaches their maximum hardness values at first and then decrease with further rising temperature. These results above indicate that the precipitation strengthening from the Cu₂NiZn precipitates has little effect on the improvement in hardness for the Cu-Zn-Ni alloy due to its small volume fraction. However, more Cu₂NiZn particles are precipitated from the matrix for the Cu-Zn-Ni-Ti alloy due to the higher dislocation density introduced during the rolling process. Furthermore, nano-sized Ni₃Ti particles are also precipitated from the matrix, resulting in better precipitation strengthening. Generally, the aging process is a competition between precipitation strengthening and the softening behavior caused by recovery and recrystallization. When the aging temperature is higher than 440 °C, over-aging occurs, and the effect of precipitation strengthening is lower than the softening behavior, thereby leading to a decrease in hardness. Figure 11b presents the variations in electrical conductivity of the samples with CR 90% deformation and aging at different temperatures for 1 h. The electrical conductivity of the Cu-Zn-Ni and Cu-Zn-Ni-Ti alloys both increase continuously with the increase in temperature. The fact illustrates that the higher aging temperature provides more nucleation energy for the precipitation of precipitates, facilitating the solute atoms to precipitate from the supersaturated solid solution during the aging process and resulting in an improvement in electrical conductivity.

Therefore, a good combination of hardness and electrical conductivity is achieved with CR 90% deformation and aging at 440 °C for 1 h for both of the Cu-Zn-Ni and Cu-Zn-Ni-Ti alloys. Thus, the analysis about the tensile properties of the alloys in Section 3.2.2 is performed at the above aging condition.

3.2.2. Tensile Properties

Figure 12 presents the typical engineering stress–strain curves of the Cu-Zn-Ni and Cu-Zn-Ni-Ti alloys after CR 90% deformation and aging at 440 °C for 1 h. The variations in yield strength (YS), ultimate tensile strength (UTS) and elongation are summarized in Table 3. In comparison to the Cu-Zn-Ni alloy, the YS and UTS of the Cu-Zn-Ni-Ti alloy exhibit remarkable improvements, whereas the elongation undergoes a slight decrease. The YS and UTS of the Cu-Zn-Ni-Ti can reach 1192 MPa and 1297 MPa, respectively, which are 229 MPa and 263 MPa higher than those of the Cu-Zn-Ni alloy. The elongation of the Cu-Zn-Ni-Ti is 1.8%, which is 1.3% lower than that of the Cu-Zn-Ni alloy.

As concluded from the above results, the addition of Ti remarkably improves the tensile strength of the Cu-Zn-Ni alloy. As a typical example, the strengthening mechanisms of Cu-Zn-Ni-Ti alloy with CR 90% deformation and aging at 440 °C for 1 h are calculated and discussed here. In fact, the yield strength increment due to different strengthening mechanisms can be expressed as [27]

(4)${\sigma }_{\mathrm{total}}={\sigma }_0+\Delta {\sigma }_{\mathrm{Orown}}+\Delta {\sigma }_{\mathrm{GB}}+\Delta {\sigma }_{\mathrm{dis}}+\Delta {\sigma }_{\mathrm{TB}}+\Delta {\sigma }_{\mathrm{SS}}$

Here, ${\sigma }_0$ is the lattice friction stress of the Cu-Zn alloy, with the value of 44 MPa [22]. $\Delta {\sigma }_{\mathrm{Orowan}}$, $\Delta {\sigma }_{\mathrm{GB}}$, $\Delta {\sigma }_{\mathrm{dis}}$, $\Delta {\sigma }_{\mathrm{TB}}$ and $\Delta {\sigma }_{\mathrm{SS}}$ are the yield strength increments contributed from Orowan strengthening, grain boundary strengthening (GBs), dislocation strengthening, twin boundary strengthening (TBs) and solid solution strengthening (SSs), respectively.

Orowan strengthening plays a significant role in improving the yield strength of the studied alloy. The yield strength increment contributed from the Orowan strengthening is estimated according to the Orowan–Ashby Equation [28]:

(5)$\Delta {\sigma }_{\mathrm{Orowan}}={\displaystyle \frac{0.81\mathrm{M}\mathrm{G}\mathrm{b}}{2\mathsf{\pi }{(1-\mathsf{\upsilon })}^{1/2}}}{\displaystyle \frac{\ln ({\mathrm{d}}_{\mathrm{pre}}/\mathrm{b})}{\lambda -{\mathrm{d}}_{\mathrm{pre}}}}$

Here, M is the Taylor factor, G is the shear modulus of the matrix, υ is the Poisson’s ratio, $d_{\mathrm{pre}}$ is the average diameter of the precipitates, and λ is the edge-to-edge inter-precipitate spacing, which is related to the volume fraction f of the second phase and can be calculated as Equation [28]:

(6)$\lambda ={\displaystyle \frac{1}{2}}{d}_{\mathrm{pre}}\sqrt{{\displaystyle \frac{3\mathsf{\pi }}{2f}}}$

Here, f is the volume fraction of Ni₃Ti and Cu₂NiZn precipitates. It is evident from the data displayed in Table 4 that the Orowan strengthening increment from Ni₃Ti precipitates is 251.2 MPa and from Cu₂NiZn precipitates is 211.7 MPa. Hence, the total yield strength contributed from Orowan strengthening is 462.9 MPa.

GBs can hinder the movement of dislocations. Generally, the yield strength increment caused by GBs is calculated by the Hall–Petch equation [33]:

(7)$\Delta {\sigma }_{\mathrm{GB}}={\displaystyle \frac{K_{H-P}}{\sqrt{d_{\mathrm{grain}}}}}$

Here, $K_{H-P}$ is the Hall–Petch coefficient for Cu-Zn alloys, and $d_{\mathrm{grain}}$ is the average diameter of the subgrains. The parameters used in the current study are presented in Table 5. Therefore, the yield strength increment from the GBs is calculated as 337.9 MPa.

As is known, the dislocation density of an alloy increases dramatically during the rolling process. The yield strength increment due to the dislocation strengthening can be calculated by the Taylor equation [30]:

(8)$\Delta {\sigma }_{\mathrm{dis}}=M\alpha G\mathrm{b}{\rho }^{1/2}$

Here, α is a geometric constant. The parameters used in the current study are listed in Table 6, and the corresponding yield strength increment due to the dislocation strengthening is calculated as 184.1 MPa.

The TBs exhibits a similar strengthening effect to that of the GBs, and the yield strength increment due to TBs is expressed by a modified Hall–Petch relationship [22]:

(9)$\Delta {\sigma }_{\mathrm{TB}}=V_{fT}{\displaystyle \frac{K_{\mathrm{TB}}}{\sqrt{d_{\mathrm{TB}}}}}$

Here, V_fT is the volume fraction of deformation twins, $K_{TB}$ is a constant, which is approximately identical to $K_{H-P}$ [22,35,36,37], and $d_{\mathrm{TB}}$ is the average twin boundary spacing. Based on the parameters given in Table 7, the yield strength increment caused by TBs is calculated as 54.4 MPa.

Typically, in comparison to the matrix, solid solution atoms have a different atomic radius and shear modulus (G). When solute atoms are introduced into the matrix, the matrix lattice becomes locally distorted. The yield strength increment due to the SSs can be estimated by the equation [22,34]:

(10)$\Delta {\sigma }_{\mathrm{SS}}=MG{\displaystyle \frac{{{\epsilon }_{SS}}^{3/2}{c}^{1/2}}{700}}$

Here, ${\epsilon }_{\mathrm{SS}}$ is the misfit strain due to a lattice distortion near the solute, and c is the concentration of solute atoms in the matrix. Assuming that 1.5 wt.% of Ti atoms completely react with Ni atoms to form the NiTi and Ni₃Ti phases, the consumption of Ni ranges between 1.84 wt.% and 5.52 wt.%, indicating that at least 0.48 wt.% of Ni atoms are estimated to dissolve into the matrix for the alloy. Based on the data list in Table 8, the yield strength increment from the SSs is calculated as ~5.2 MPa.

Therefore, the total calculated value of yield strength of the Cu-Zn-Ni-Ti alloy is 1088.5 MPa, which agrees well with the experimental result (1192 MPa), with an error of only 8.7%. This slight difference may occur due to the error in estimating the size of precipitates, deformation twins width, dislocation density and volume fraction of the precipitates and deformation twins. The contribution ratios of each strengthening mechanism to the yield strength of the alloy are presented in Figure 13. It is noticeable that the Orowan strengthening is the primary strengthening mechanism of the alloy, which accounts for 42.5% of the total yield strength, and then the GBs (31.1%), the dislocation strengthening (16.9%), the TBs (5.0%) and the SSs (0.5%), accordingly.

3.3. Corrosion Behavior

3.3.1. Average Corrosion Rate

The samples with the best combination of mechanical properties and electrical conductivity (CR 90% deformation and aging at 440 °C for 1 h) were selected for corrosion tests. The average corrosion rates of the aged Cu-Zn-Ni and Cu-Zn-Ni-Ti alloys in 3.5% NaCl solution for 30 days are calculated according to the ASTM standard G31-72(2004) [39]:

(11)$R={\displaystyle \frac{250\times \mathrm{K}\times W}{3\times A\times T\times D}}$

Here, R is the corrosion rate in μm/month; K is a constant, with the value of 8.76×10⁴; W is the mass loss (g); A is the exposed surface area (cm²); T is the exposure time (h); and D is the density ($\mathrm{g}/{\mathrm{cm}}^3$) of the samples. The variations in W, A and D of the two alloys are listed in Table 9. The average corrosion rates of the Cu-Zn-Ni-Ti and Cu-Zn-Ni alloys are measured as 2.8 μm/month and 4.3 μm/month, respectively, indicating that the Cu-Zn-Ni-Ti alloy exhibits a better corrosion resistance.

3.3.2. Potentiodynamic Polarization

Figure 14 displays the potentiodynamic polarization curves of the Cu-Zn-Ni and Cu-Zn-Ni-Ti alloys in a 3.5 wt.% NaCl solution. The data of E_corr, I_corr, ${\beta }_{\mathrm{a}}$ and ${\beta }_{\mathrm{c}}$ obtained from the Tafel plots are presented in Table 10. There is no significant difference in E_corr between the Cu-Zn-Ni and the Cu-Zn-Ni-Ti alloy, but the I_corr of the Cu-Zn-Ni alloy decreases from $2.0\times {10}^{-6}$ A·cm⁻² to $5.1\times {10}^{-7}$ A·cm⁻² for the Cu-Zn-Ni-Ti alloy. Such a result shows clearly that Ti addition can effectively improve the corrosion resistance of the Cu-Zn-Ni alloy.

3.3.3. Corroded Surface Morphology

Figure 15 exhibits the corroded surface morphologies of the aged Cu-Zn-Ni and Cu-Zn-Ni-Ti alloys after being exposed to the 3.5% NaCl solution for 30 days. It is noticeable from Figure 15a that the Cu-Zn-Ni alloy surface is severely attacked by the corrosive medium, and a large amount of pitting is formed in the matrix. Additionally, a small amount of corrosion product is attached to the matrix surface, and micro-cracks appear in the corrosion products. In comparison to the Cu-Zn-Ni alloy, the amount of pitting is significantly decreased with Ti addition (Figure 15b), illustrating that the Cu-Zn-Ni-Ti alloy has superior corrosion resistance compared with the Cu-Zn-Ni alloy. The results above imply that Ti is preferentially corroded and forms a dense TiO₂ layer on the surface [17], which hinders the anodic dissolution of Ni and Zn in the matrix and alleviates the deterioration of corrosion resistance.

It can be inferred that the addition of Ti is beneficial to improve the hardness, tensile strength, electrical conductivity and corrosion resistance of the Cu-Zn-Ni alloy. After suitable thermo-mechanical treatments, a desired combination of hardness (327.1 HV), YS (1192 MPa), UTS (1297 MPa), elongation (1.8%), electrical conductivity (11.2% IACS) and corrosion rate (2.8 μm/month) are obtained for the Cu-30Zn-6Ni-1.5Ti alloy. More comprehensive works on the further content optimization, precipitate phase transformation behavior and corrosion mechanisms of the alloy will be presented in future studies.

4. Conclusions

In the present study, the effects of Ti addition on microstructure, mechanical properties and corrosion resistance of the Cu-Zn-Ni alloy have been investigated in detail, and the main conclusions are summarized below:

• (1). Microstructure analysis confirms that the Ni₃Ti and NiTi phases are formed in the Cu-Zn-Ni-Ti alloy. Most of the Ni₃Ti particles dissolve into the matrix, whereas NiTi particles remain after the solution treatment. Moreover, nano-sized Ni₃Ti and Cu₂NiZn phases are precipitated from the matrix during the aging process.

• (2). The yield strength improvement of the studied alloy are attributed to the Orowan strengthening (accounting for 42.5% of the total yield strength), then the grain boundary strengthening (31.1%), the dislocation strengthening (16.9%), the twin boundary strengthening (5.0%) and the solid solution strengthening (0.5%).

• (3). After CR 90% deformation and aging at 440 °C for 1 h, the designed alloy has a hardness of 327.1 HV, a YS of 1192 MPa, a UTS of 1297 MPa, an elongation of 1.8%, an electrical conductivity of 11.2% IACS and a corrosion rate of 2.8 μm/month in the 3.5% NaCl solution.

Author Contributions

Conceptualization, J.J. and T.L.; methodology, T.L.; software, J.J.; validation, Z.Y., H.C. and J.L.; formal analysis, X.S.; investigation, H.C.; resources, Y.W.; data curation, X.S.; writing—original draft preparation, X.S.; writing—review and editing, J.J.; visualization, J.L.; supervision, Z.Y.; project administration, Y.W.; funding acquisition, Y.W. All authors have read and agreed to the published version of the manuscript.

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.

Footnotes

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Figures and Tables

Figure 1. The relation between peak broadening βcos(θ) and 4sin(θ) for the Cu-Zn-Ni-Ti alloy after CR 90% deformation and aging at 440 °C for 1 h.

Figure 2. The dimension of the tensile specimens.

Figure 3. Optical micrographs of the as-cast (a) Cu-Zn-Ni and (b) Cu-Zn-Ni-Ti alloys.

Figure 4. SEM images of the as-cast (a) Cu-Zn-Ni, and (b) Cu-Zn-Ni-Ti alloys.

Figure 5. (a) XRD patterns of the as-cast Cu-Zn-Ni and Cu-Zn-Ni-Ti alloys and (b) magnified patterns of the area in the rectangle box in (a) between 34° and 52°.

Figure 6. Microstructures of the (a,c) Cu-Zn-Ni and (b,d) Cu-Zn-Ni-Ti alloys after solution treatment.

Figure 7. SEM images of the (a) Cu-Zn-Ni and (b) Cu-Zn-Ni-Ti alloys with CR 90% deformation.

Figure 8. TEM images of the Cu-Zn-Ni alloy with CR 90% deformation and aging at 440 °C for 1 h: (a) dislocation cells, (b) Cu2NiZn precipitates and corresponding FFT image, (c) subgrains and (d) twins.

Figure 9. TEM images of the Cu-Zn-Ni-Ti alloy with CR 90% deformation and aging at 440 °C for 1 h: (a) NiTi phase and subgrain, (b) SAED pattern of the NiTi phase, (c) Ni3Ti precipitates and dislocation cells, (d) corresponding FFT image of Ni3Ti precipitates, (e) Cu2NiZn precipitates and twins and (f) corresponding FFT image of Cu2NiZn precipitates.

Figure 10. Variations in (a) hardness, and (b) electrical conductivity of the Cu-Zn-Ni and Cu-Zn-Ni-Ti alloys with CR 90% deformation and aging at 440 °C for various times.

Figure 11. Variations in (a) hardness and (b) electrical conductivity of the Cu-Zn-Ni and Cu-Zn-Ni-Ti alloys with CR 90% deformation and aging at different temperatures for 1 h.

Figure 12. Stress–strain curves of the Cu-Zn-Ni and Cu-Zn-Ni-Ti alloys after CR 90% deformation and aging at 440 °C for 1 h.

Figure 13. The contribution ratio of each strengthening mechanism to the yield strength of the Cu-Zn-Ni-Ti alloy with CR 90% deformation and aging at 440 °C for 1 h.

Figure 14. Potentiodynamic polarization curve tests of the Cu-Zn-Ni and Cu-Zn-Ni-Ti alloys.

Figure 15. Corroded surface morphologies (a,b) of the aged Cu-Zn-Ni and Cu-Zn-Ni-Ti alloys after being exposed to 3.5% NaCl solution for 30 days.

Table 1

Chemical composition of the alloys determined by XRF-1800.

Alloy	Cu (wt.%)	Zn (wt.%)	Ni (wt.%)	Ti (wt.%)
Cu-30Zn-6Ni	Bal.	29.13	6.06	-
Cu-30Zn-6Ni-1.5Ti	Bal.	29.75	5.98	1.43

Table 2

EDS results of the alloys at different positions.

Alloy	Point	Chemical Composition, at.%
		Cu	Zn	Ni	Ti
	A	5.01	3.98	91.01	-
Cu-Zn-Ni	B	3.82	4.79	91.39	-
	C	62.56	31.95	5.49	-
Cu-Zn-Ni-Ti	D	2.32	3.47	46.73	47.48
	E	35.56	15.75	37.16	11.53
	F	60.49	26.73	9.72	3.06
	G	59.58	33.87	5.63	0.92

Table 3

Variations in mechanical properties of the Cu-Zn-Ni and Cu-Zn-Ni-Ti alloys after CR 90% deformation and aging at 440 °C for 1 h.

Alloy	Young’s Moduli (GPa)	YS (MPa)	UTS (MPa)	Elongation (%)
Cu-Zn-Ni	131	963 ± 25	1034 ± 32	3.1 ± 0.5
Cu-Zn-Ni-Ti	133	1192 ± 34	1297 ± 41	1.8 ± 0.3

Table 4

The parameter values used in yield strength calculation of the alloy contributed from the Orowan strengthening.

Parameter	Value	Units	Refs.
M	3.06	-	[29]
G	41	GPa	[30]
b	0.2556	nm	[31]
v	0.33	-	[32]
$d_{{\mathrm{Ni}}_3\mathrm{Ti}}$	15	nm	This paper
$d_{{\mathrm{Cu}}_2\mathrm{NiZn}}$	5	nm	This paper
$f_{{\mathrm{Ni}}_3\mathrm{Ti}}$	2.82	%	This paper
$f_{{\mathrm{Cu}}_2\mathrm{NiZn}}$	0.51	%	This paper

Table 5

The parameter values used in yield strength calculation of the alloy contributed from the grain boundary strengthening.

Parameter	Value	Units	Refs.
$K_{H-P}$	0.11	$\mathrm{MPa}·{\mathrm{m}}^{1/2}$	[22]
$d_{\mathrm{grain}}$	106	nm	This paper

Table 6

The parameter values used in yield strength calculation of the alloy contributed from the dislocation strengthening.

Parameter	Value	Units	Refs.
α	0.2	-	[34]
ρ	$8.24\times {10}^{14}{\mathrm{m}}^{-2}$	m−2	This paper

Table 7

The parameter values used in yield strength calculation of the alloy contributed from the twin boundary strengthening.

Parameter	Value	Units	Refs.
VfT	8.7	%	This paper
d TB	31	nm	This paper

Table 8

The parameter values used in yield strength calculation of the alloy contributed from the solid solution strengthening.

Parameter	Value	Units	Refs.
${\epsilon }_{SS}$	0.56	${\mathrm{MPa} \mathrm{m}}^{1/2}$	[38]
c	0.48	wt.%	This paper

Table 9

Parameters used in corrosion rate calculation.

Alloy	W (g)	A (cm2)	D (g/cm3)	Corrosion Rate(μm/month)
Cu-Zn-Ni	$(4.1\pm 0.1)\times {10}^{-2}$	11.7 ± 0.5	8.5 ± 0.1	4.3 ± 0.4
Cu-Zn-Ni-Ti	$(2.9\pm 0.1)\times {10}^{-2}$	12.1 ± 0.6	8.4 ± 0.1	2.8 ± 0.3

Table 10

Electrochemical parameters of Cu-Zn-Ni and Cu-Zn-Ni-Ti alloys derived from Tafel plots.

Alloy	Ecorr (V)	Icorr (A·cm−2)	${\mathit{\beta }}_{\mathbf{a}}$ (V−1)	${\mathit{\beta }}_{\mathbf{c}}$ (V−1)
Cu-Zn-Ni	−0.28 ± 0.04	$(2.0\pm 0.1)\times {10}^{-6}$	0.13 ± 0.01	0.17 ± 0.01
Cu-Zn-Ni-Ti	−0.26 ± 0.03	$(5.1\pm 0.3)\times {10}^{-7}$	0.20 ± 0.02	0.09 ± 0.01

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