Nickel-aluminum bronze (NAB) alloys are widely used for marine propellers, pumps, and valves due to the attractive mechanical properties and excellent corrosion resistance [[1], [2], [3]]. The microstructure of as-cast NAB alloys is generally composed of α phase, martensite β phase (β' phase) and κ phases, of which κ_I, κ_II, κ_IV phases (based on Fe₃Al) are rich in iron and κ_III phase (based on NiAl) is rich in nickel [[4], [5], [6]]. As-cast NAB alloys often suffer from defects, composition segregation, and inhomogeneous structure, which are detrimental to the mechanical properties and corrosion resistance, and thus limiting their applications.

It is well known that the alloy properties are closely related to the microstructure [[7], [8], [9], [10], [11]]. Recently, two methods for improving the microstructure and properties of NAB alloys have been developed. One is to modify the microstructure and properties of the NAB alloy substrate via composition design [[12], [13], [14]], plastic deformation techniques (equal channel angular pressing (ECAP) [15], hot rolling [16], etc.) and heat treatment [17,18]. Yang et al. [14] found that the increase of Ni content in NAB alloy could reduce the content of highly corrosive β' phase in the microstructure and make the surface passive film more compact, thereby enhancing corrosion resistance. Wharton et al. [19] studied the corrosion properties of as-cast and as-forged NAB alloy using a series of electrochemical techniques. They found that the corrosion rate of NAB alloy was mainly related to the characteristics of surface protective film, alloy composition, heat treatment process, and composition of corrosive medium. Lv et al. [20] found that more β phase transformation products were formed in the forged NAB alloy during the rolling process, and the mechanical properties were significantly enhanced. The other one is to modify the surface of materials directly in contact with the external environment, such as friction stir processing (FSP) [[21], [22], [23], [24]], laser surface melting [25,26] and ion implantation [27,28]. For example, Ni et al. [21,29] reported that the surface microstructure of the NAB alloy treated by FSP was finer than that of the as-cast alloy, and the strength, hardness, and plasticity were significantly improved. Moreover, the NAB alloy treated by FSP exhibited better corrosion resistance in a neutral chloride solution. Nickel ion implantation was found to effectively suppress selective phase corrosion and improve the corrosion resistance of the NAB alloy [27]. A denser film formed on the surface of as-cast NAB alloy after ion implantation was responsible for the improved properties.

Studies have shown that corrosion mainly occurs at the eutectoid structures in the NAB alloys [30,31]. Song et al. [32] found that the interface between the lamellar phase and corrosion product is the channel, through which the corrosive medium enters the matrix of nickel-aluminum bronze. Therefore, the corrosion resistance of the alloy can be optimized by improving the morphology of the eutectoid structures. Furthermore, it is reported that the lamellar microstructures of nickel-aluminum bronze could be spheroidized by ECAP at 400 °C and 600 °C [15]. Hot rolling, similar to ECAP, can also produce large deformation at high temperature, which may also spheroidize the continuous lamellar phase [33,34] and make the corrosion channel become discontinuous. In addition, hot rolling can eliminate defects and make the microstructure more homogeneous, which also improves the corrosion resistance of the alloy. Although there are a few studies reporting the effect of hot rolling on the microstructure and mechanical properties of nickel-aluminum bronze [16,20], there are few reports on the use of hot rolling to enhance the corrosion resistance of nickel aluminum bronze, and thus the related corrosion mechanisms have not been revealed. Therefore, it is necessary to study the effect of hot rolling on the properties of nickel-aluminum bronze, especially on the corrosion resistance.

In this study, a high-nickel content as-cast nickel-aluminum bronze alloy (Cu-9Al-10Ni-4Fe-1.2Mn, all in wt.%) is used as the experimental material. The effects of hot rolling on the microstructures, mechanical properties, and corrosion resistance of NAB alloys are studied in this work. The corrosion properties are evaluated by mass loss tests, electrochemical impedance spectroscopy (EIS) tests, and cross-sectional morphology examinations. The relationship between the microstructure of NAB alloy and mechanical properties, as well as corrosion resistance, is the focus of this work, and thus will be discussed in detail.

The as-cast NAB alloy (with the nominal composition of Cu-9Al-10Ni-4Fe-1.2 Mn) studied in this work was melted in a vacuum induction melting furnace under argon atmosphere. The melt was cast into an iron mould with a size of 45 mm × 45 mm × 150 mm. The real composition (wt.%) of the as-cast ingot, as measured by X-ray fluorescence spectrometry method, is 8.92 % Al, 9.79 % Ni, 3.91 % Fe, 1.11 % Mn, and balance Cu. A part of the as-cast ingot was annealed at 675 °C for 6 h and then cooled in the furnace (termed as as-annealed sample).

The samples with a size of 20 mm × 20 mm × 88 mm were machined from the as-annealed sample and then heated and equilibrated at the intended rolling temperature (850 °C) for 20 min. The hot rolling operation was carried out using a laboratory rolling mill. A reduction thickness of ~1.5 mm for each successive pass (the reduction thickness of the last pass was 1 mm, 11 passes in total) was used. The sample temperature was maintained at 850 °C for 20 min between each successive pass. After the final pass, the samples were air-cooled to room temperature. This resulted in the total deformation of 80 % and the final thickness of 4 mm for each sample. For comparison, a part of the as-annealed NAB sample was further held at 850 °C for 220 min, followed by air cooling.

After being mechanically ground and polished, the as-cast, as-annealed, and hot-rolled samples were etched for 10 s, with a solution containing 3 g FeCl₃, 2 mL HCl (36-38 wt.%) and 95 mL C₂H₅OH. The microstructures and chemical compositions were studied by using an optical microscope (OM, Olympus GX51, Olympus Corp., Japan) and a scanning electron microscope (SEM, Zeiss supra 55, Zeiss Corp., Germany) equipped with an energy dispersive spectroscopy (EDS).

The samples for transmission electron microscope (TEM) investigation were ground to 40 μm initially and then ion-beam milled using a Gatan precision ion polishing system (Gatan model 691, Gatan Inc., America). The TEM foils were characterized using a field emission transmission electron microscope (JEM 2100 F, JEOL Ltd., Japan).

Samples for tensile test were machined from the as-cast, as-annealed, and hot-rolled samples (the sample axis is parallel to the rolling direction). The samples were tested on a universal tensile machine (Instron 5569, Instron Corp., America) with a loading rate of 2 mm/min at room temperature (~25 °C). At least three samples were tested for each condition. Tensile fracture morphology was observed by SEM.

Vickers hardness was measured by a micro-hardness tester (MH-50, Shanghai Everone Precision Instruments Corp., China) with an applied load of 500 g and a holding time of 15 s. For each sample, twelve indentations were conducted at different positions to obtain the average value.

Samples for static immersion test (a dimension of 30 mm × 15 mm × 2 mm, with a circular hole of 3 mm in diameter) were mechanically ground with 240 grit SiC abrasive paper. All samples were treated by ultrasonic cleaning in ethanol for 3 min, rinsed by distilled water, dried in cold air, and then weighed carefully on an analytical balance (ME204E, Mettler Toledo Corp., Switzerland). The exposed area of each sample was calculated independently. Each test sample was immersed into a container containing 300 mL of 3.5 wt.% NaCl solution under quiescent condition at room temperature (25 ± 1 °C). The volume of the solution exceeded the minimum solution volume to sample area ratio (0.20 mL/mm²) required by ASTM Standard G31-72 [35]. The NaCl solution was prepared from distilled water and analytical grade chemicals. During the immersion test, the solution was replaced with a fresh one at 7-day intervals. After being immersed for 1, 3, 5, 7, 14, and 28 days, respectively, the corroded samples were taken out and immersed in the solution of 500 mL H₂O + 500 mL HCl (36-38 wt.%, AR) for 3 min to clean off the surface corrosion products. Then these samples were completely rinsed with distilled water, dried in a thermostatic drying chamber at 60 °C for 2 h, and finally weighed. Three samples were tested for each condition, and an average value of mass loss was obtained.

Another group of samples for corrosion test were ground with SiC paper down to 2000 grit and polished with 1 μm diamond paste, and then immersed in 3.5 wt.% NaCl solution. For the surface morphology observation and phase identification, the samples were immersed for 1, 3, 30 days, respectively, and then the surface morphology of each sample was observed by SEM. The phases present in the corrosion product film were analyzed by using an X-ray diffractometer (XRD, PANalytical Empyrean, Netherlands) using Cu K_α radiation, working at 40 mA and 40 kV. For the morphology observation at the cross-section, the samples were immersed for 60 days and characterized by using an electron probe microscopic analyzer (EPMA, JXA-8530 F PLUS, JEOL Ltd., Japan).

The electrochemical corrosion studies were performed in naturally aerated 3.5 wt.% NaCl solution at room temperature (25 ± 1 °C) using a Gamry electrochemical workstation (Reference 600, Gamry Instruments Inc., America). A three-electrode system, where the studied alloy was taken as working electrode (WE), a large platinum foil as counter electrode (CE) and a Ag/AgCl electrode (saturated KCl with electrode potential of 0.1981 V vs. Standard Hydrogen Electrode) as reference electrode (RE), was employed.

The samples were mechanically ground with SiC abrasive paper starting from 240 to 2000 grit size, washed with distilled water, dried by cold air and then transferred quickly to the 3.5 wt.% NaCl solution. The working area of each sample exposed to tested solution was 10 mm in diameter and 0.7854 cm² in area. Each experiment was carried out in a freshly prepared solution.

After the open circuit potential (OCP) was steady, the electrochemical impedance spectroscopy (EIS) measurement was carried out at OCP in the frequency range from 100,000 Hz to 0.01 Hz. To ensure good reproducibility, at least three samples were tested for each condition. According to the general variation trend of experimental results, the most representative EIS data were selected for fitting using Zsimpwin 3.10 software package.

The microstructures of the NAB samples under different processing conditions were illustrated in Fig. 1, Fig. 2, Fig. 3, respectively. As shown in Fig. 1(a), the OM microstructure of the as-cast NAB sample was composed of large blocky α phase (light etched areas), as indicated by the a1 area in Fig. 1(a), and intermetallic κ precipitates (dark etched areas). The details of α phase and intermetallic κ precipitates can be further observed in the SEM image (Fig. 2(a)). According to the EDS results in Table 1, the α phase (Point A) is a copper-rich solid solution; lamellar κ_III (Point B) is rich in Ni [14], and the coarse IMC strip (Point C), which is rich in Ni and Al, has been ascertained to be NiAl phase [14].

It is obvious that the size of the eutectoid structure in the as-annealed NAB sample was increased (Fig. 1(b)), and a large number of small phases were precipitated in α phase (the b1 area in Fig. 1(b)). As shown in Fig. 2(b), the lath-shaped phase (Point F) precipitated in the α phase was densely distributed and rich in Ni and Al (see Table 1). The κ_III phases in the α + κ_III eutectoid appear as lamellar structure [6]. The distribution between the lamellar κ_III phases and lamellar α phases are similar to parallel arrangement, as shown in the SEM image in Fig. 2(b) and TEM image in Fig. 3(a). In addition, the lamellar κ_III phases are much longer than the lath-shaped phases precipitated in α phase. For example, The length of κ_III phase in Fig. 3(a) is more than 3730 nm, and can be identified as NiAl phase [6] according to the selected area diffraction pattern (SADP, see the inset in Fig. 3(a)). The SADP (inset in Fig. 3(b)) also confirms that the lath-shaped phase (rich in Ni and Al according to Table 1) precipitated in α phase is NiAl based κ_V phase [2,36]. κ_V phases which distributed disorderly in α matrix, are much finer than the lamellar κ_III phases. The length of the lath-shaped κ_V phase in Fig. 3(b) is only 1412 nm.

In the hot-rolled NAB sample, as shown in Fig. 1(d), the original lamellar phases were fragmented and partially spheroidized. The secondary phases within α grains were completely spheroidized and uniformly distributed in the α matrix (the d1 area), and the α grains became slender along the rolling direction. It can be observed from the SEM image in Fig. 2(c) that the original lamellar κ_III was fragmented (Point H), and that the precipitates in α phase were globalized (Point I, rich in Ni and Al according to Table 1). The partially spheroidized κ_III phase (see Fig. 3(c)) and the globalized κ_V phase (Fig. 3(d)) can be identified as the NiAl phase. Moreover, dislocations and twins were generated during the hot rolling process, as shown in Fig. 3(e).

For comparison, the microstructure of the as-annealed NAB sample further held at 850 °C for 220 min is shown in Fig. 1(c). It can be found that only a few consecutive lamellar phases were slightly dissolved, and the number of precipitated phases in the α matrix was reduced (the c1 area).

Fig. 4 presents the typical tensile stress-strain curves and the corresponding tensile fracture morphology of NAB samples under different conditions. The ultimate tensile strength, yield strength, and elongation are summarized in Table 2, where the micro-hardness of each NAB sample is also included.

As shown in Table 2, the as-cast NAB sample exhibits the lowest strength and hardness, while the highest elongation. After annealing, the ultimate tensile strength, yield strength, and hardness of the NAB sample increased. However, the elongation decreased from 18.6 ± 1.2% to 11.4 ± 0.7 %. After hot rolling at 850 °C, the ultimate tensile strength (882.8 ± 5.5 MPa) and yield strength (658.7 ± 4.9 MPa) of the NAB sample were increased by ~28.2 % and ~84.2 %, respectively, compared to that of the as-annealed NAB sample. Moreover, the elongation increased from 11.4 ± 0.7% to 15.2 ± 0.5 %.

Fig. 4(b) shows the fractography of the as-cast sample after the tensile test. It can be observed in the as-cast sample that the fracture morphology exhibits ductile deformation, as indicated by both dimples (the circular mark) and tearing features. The as-annealed sample exhibits river-like morphology (Fig. 4(c)), which is slender and finer than that of the as-cast NAB sample. After hot rolling at 850 °C, the sample exhibits a ductile fracture feature with the presence of a great number of dimples with various sizes (Fig. 4(d)). The slender α grains were fractured after plastic deformation and the precipitates were retained in the fractured α grains.

3.3.1. Static immersion test

Fig. 5(a) shows the variation of the mass loss of the NAB samples with the immersion time. It can be seen that the mass loss of the three samples gradually increased with the increase of immersion time. After annealing, the mass loss of the NAB sample increased significantly. For example, the mass loss increased from 3.53 ± 0.08 g/m² to 14.38 ± 0.31 g/m² after 28 days immersion. However, after hot rolling, the mass loss of the NAB sample decreased dramatically to 1.86 ± 0.09 g/m² after immersion for 28 days, which was decreased by ~87 % compared to the as-annealed sample.

Corrosion rate can be calculated according to ASTM Standard G31-72(2004) [35]:

where R is the corrosion rate in mm/year, K the constant (8.76 × 10⁴), W the mass loss (g), A the exposed area of the sample (cm²), T the immersion time (h), and D the density of sample (g/cm³). The corrosion rate curves are shown in Fig. 5(b). The corrosion rate of all the samples in the initial stage of immersion (0-3 days) dropped sharply, decreased gradually in 3-14 days, and finally came to an even lower value after 14 days. The corrosion rate of the as-annealed NAB sample was increased compared to that of the as-cast sample. After hot rolling, the corrosion rate was greatly reduced to 0.0032 ± 0.0001 mm/year after immersion for 28 days.

3.3.2. Electrochemical impedance measurement

The samples were initially immersed in 3.5 wt.% NaCl solution for 1 day and 5 days, respectively, and then subjected to electrochemical impedance spectroscopy (EIS) test. Nyquist plots of these samples are presented in Fig. 6. The symbols and solid lines represent the measured data and fitted curves, respectively.

The equivalent electrical circuit model [14,27,28,37,38] commonly used to explain the corrosion behavior of NAB samples in chloride solutions is shown in Fig. 6(c). Specifically, R_s represents the solution resistance; CPE is the constant phase element; R_f and CPE₁ correspond to the resistance and non-ideal capacitance of the passive film formed on the surface of the sample, respectively; R_ct and CPE₂ represent the charge transfer resistance and the non-ideal capacitance for double-charge layer, respectively. W_d is the Warburg diffusion impedance. The impedance of CPE can be defined as follows [39]:

where Q is the magnitude of the CPE, j is the imaginary root, ω is the angular frequency, and n is the CPE exponent [39,40]. The polarization resistance R_p can be expressed as R_p = R_f + R_ct. The fitting results are listed in Table 3. In all cases, the chi-square value is in the order of 10^-3 magnitude, indicating that the fitting for the data is acceptable.

As shown in Fig. 6(a), the Nyquist plots of the three samples after 1 day immersion in 3.5 wt.% NaCl solution exhibits the similar shape in the higher frequency (HF) range with a capacitive reactance arc. The capacitive arc in the higher frequency range represents a combination of charge transfer resistance on bare electrode and that contributed by the passive film [41]. As the immersion time extended from 1 day to 5 days (Fig. 6(b)), the diameter of HF capacitive arc of all the samples increased. Besides, the diameter of HF capacitive arc of as-annealed sample was reduced compared to the as-cast one (Fig. 6(a) and (b)) after the same immersion time; the film resistance (R_f) and the polarization resistance (R_p) were also greatly reduced (Table 3). However, the diameter of HF capacitive arc of the hot-rolled sample was greatly increased and was larger than that of the as-cast sample. According to Table 3, the R_f and R_p of the hot-rolled sample after immersing for 5 days were increased by 375 % and 385 %, respectively, compared to that of the original as-cast sample, indicating that the corrosion resistance of the NAB samples was improved by hot rolling.

3.3.3. Corrosion products and corrosion morphology

Fig. 7 shows the XRD patterns of the NAB samples after immersion in 3.5 wt.% NaCl solution for 30 days. Copper chloride hydroxides were identified in all the three NAB samples but had the lowest intensity in the hot-rolled NAB sample. Cu₂O was only detected in the as-annealed sample and was not identified in the other two samples, which was probably due to the small amount of Cu₂O in the as-cast and hot-rolled samples. The intensity of NiAl peak was increased significantly after annealing, while decreased slightly after further hot rolling. This result is following the observation of microstructures in section 3.1.

Fig. 8 shows the surface morphology of the samples after immersion for different time in 3.5 wt.% NaCl solution. After immersing for 1 day, it can be seen that corrosion mainly occurred at the lamellar α + κ_III phases for the as-cast NAB sample (Fig. 8(a2)), which is consistent with previous studies [30,31]. With the increase of immersion time, the surface corrosion of the sample became severe. A number of blocky corrosion products (indicated by Point A) appeared on the surface after 30 days of immersion (Fig. 8(a4)). The chemical composition (wt.%) of the blocky corrosion products is 1.80 % Cl, 28.88 % O, 51.71 % Cu, 12.43 % Al, 2.98 % Fe, 2.21 % Ni, indicating that the corrosion products are the copper chloride hydroxides [42]. The corrosion of the as-annealed NAB sample was more severe than the as-cast sample, as shown in Fig. 8(b2)-(b4). Under the same immersion condition, more corrosion products appeared on the surface of the as-annealed NAB sample (Fig. 8(b4)). The EDS analysis shows that the composition of area B is 25.24 % O, 51.51 % Cu, 9.85 % Cl, 5.70 % Ni, 4.61 % Al, 2.66 % Fe, 0.42 % Mn (all in wt.%). The high contents of Cl and O suggest that these corrosion products are the copper chloride hydroxides [42], which is also supported by the XRD results in Fig. 7. In contrast, the surface of the hot-rolled sample was corroded lightly. Only a few white corrosion products can be observed on its surface (Fig. 8(c4)).

Fig. 9 shows the back-scattered electron (BSE) images of the samples in cross-section after immersion in neutral 3.5 wt.% NaCl solution for 60 days. As shown in Fig. 9(a), the eutectoid α phase in the as-cast sample was preferentially corroded, and the maximum local corrosion depth was ~4.9 μm. The lamellar κ_III was retained in the corrosion product, and the α matrix, which is not adjacent to κ phases, was only slightly eroded. As shown in Fig. 9(b), a large number of corrosion products formed on the surface of the as-annealed NAB sample. The lamellar κ_III phase and a great number of lath-shaped κ_V phases were retained in the corrosion products. The α matrix was severely corroded with a maximum local corrosion depth of ~15.8 μm. However, the corrosion of the hot-rolled samples (Fig. 9(c)) was significantly alleviated, and the maximum local corrosion depth was only ~2.3 μm.

In this study, the as-cast NAB sample contains a high content of Ni. According to Yang et al. [14], the Ni element in the as-cast sample partially precipitates as the κ_III phase, and the remaining Ni dissolves in the α matrix. It can be found from Fig. 2 and Table 1 that the Ni content of the α phase is decreased after annealing, and a large number of Ni-rich κ_V phases are precipitated, suggesting that part of the dissolved Ni in the α phase is precipitated in the form of κ_V phase during the annealing process. Besides, the eutectoid structure is coarsened due to the accelerated element diffusion at high temperature. After hot rolling, the κ_V phase in the α matrix is spheroidized, and the lamellar κ_III phase is partially spheroidized as well. In contrast, most of the κ phases in the microstructure of the sample incubated at 850 °C for 220 min does not spheroidize, suggesting that the spheroidization of the κ phases requires the combination of high temperature and rolling. The lamellar κ_III phase and lath-shaped κ_V phase are broken by the stress imposed on the phases and spheroidized by atom diffusion. Barr et al. [43] reported that severe buckling produced by ECAP could cause the spheroidization of the lamellar κ_III phase. Therefore, both the current study and reported research show that severe deformation should be an important factor that causes the spheroidization of the lamellar κ_III phase. In addition, the twins are also found in the hot-rolled sample, as shown in Fig. 3(e). The high-density dislocations, as indicated by red arrows, are observed inside the twins. The dislocation walls subdivide the lamellar twins into several parts, which are typical characteristics of deformation twins [44,45].

It is well known that the mechanical properties are closely related to the microstructure of alloys. For the NAB samples studied in this work, the α phase is a Cu-rich solid solution with the lowest hardness and highest ductility. The κ precipitates contain a series of intermetallic compounds and are stronger than the α phase, but less ductile [46]. The as-cast NAB sample shows the highest ductility and the lowest strength due to a large amount of soft α phase (there are no hard precipitates within the α grains), and fewer κ_III phase. After annealing, the α+κ_III eutectoid structure is coarsened and a large number of lath-shaped к_V phase is precipitated from the α phase. Thus, the α matrix is strengthened by the hard κ phases (κ_III and κ_V), leading to the increase in ultimate tensile strength, yield strength, and hardness, and the decrease in ductility/elongation of the as-annealed NAB sample. After hot rolling, κ phases are fragmented, spheroidized, and uniformly distributed in the α matrix, which causes a strengthening effect on the α matrix. Moreover, the generation of dislocations during the rolling process (Fig. 3(e)) also strengthens the alloys.

Furthermore, when the lamellar or lath-shaped precipitates are subjected to external force, a large stress concentration is generated, and the micro-cracks are easy to form. However, the stress concentration caused by the spherical precipitates is relatively small, and thus the sample exhibits better plasticity. As shown in this work, the lamellar κ_III or lath-shaped κ_V phase in the annealed microstructure is partially or fully spheroidized after rolling, resulting in the improved elongation of the hot-rolled NAB sample.

As indicated by the white arrow in Fig. 4(b), in the as-cast NAB sample, the cracks generally propagate along the boundaries between the bulk α phase and the hard precipitates [47], such as κ phases and coarse NiAl strips (as indicated by the white arrows in Fig. 1(a)). The α phase possesses good plasticity since there are no hard precipitates within the soft α phases (the a1 area in Fig. 1(a)). The soft α phase mainly contributes to the excellent ductility of the NAB alloys [14,46]. After a long period of deformation, the fracture morphology shown in Fig. 4(b) is finally formed. The dimples can also be observed from Fig. 4(b) (as indicated by the circular mark). Moreover, the content of the soft α phase in the as-cast NAB sample is the highest among all the samples. Therefore, the as-cast NAB alloy exhibits the best ductility (18.6 %). As shown in Fig. 4(c), the content of κ phases increases significantly in the as-annealed sample. The cracks are likely to propagate along with the interfaces between the κ phases and the α matrix in the as-annealed sample, so that the slender river-like morphology is formed. Xu et al. [47] also reported that the fatigue crack prefers to propagate through the α+κ_III lamellae in the as-annealed NAB sample. After hot rolling, the κ phases are spheroidized and the cracks propagate along the globular phase boundary. Therefore, a great number of dimples with various sizes (Fig. 4(d)) are formed.

At present, the dissolution of Cu is recognized as the major anodic reaction during the corrosion process of NAB alloy in neutral NaCl solution [[48], [49], [50], [51]]:

The cathodic reaction is mainly the oxygen reduction reaction [[52], [53], [54], [55]]:

Then, a cuprous oxide is formed based on the following reaction [56,57]:

As the immersion time increases, Cu₂O is further oxidized to form copper chloride hydroxides, such as Cu₂(OH)₃Cl [32,37]:

At the onset of the static immersion test, a protective film rapidly forms on the surface of the NAB alloy in the 3.5 wt.% NaCl solution, and thus the corrosion rate of the samples decreases sharply. With the prolonged immersion time, the corrosion rate finally comes to an even lower value due to the continuous formation of a stable oxide film. It was reported that the corrosion rates of the NAB samples decrease with the prolonged immersion time, and finally tends to be steady [14,21]. Besides, the hot-rolled sample shows the best corrosion resistance among the three samples at the same immersion time. Comparing the EIS results of 1 day and 5 days immersion, the film resistance (R_f) and polarization resistance (R_p) of the NAB samples increase with the prolonged immersion time, reflecting the growth of the surface passive film and the improvement of the protecting effect on the substrate [58]. In addition, the R_p (the overall corrosion resistance) of the hot-rolled NAB sample is significantly higher than that of the as-cast and as-annealed one under the same immersion time, suggesting that hot rolling significantly improves the corrosion resistance of the NAB sample, which is in good agreement with the mass loss results.

Fig. 10 shows schematically the corrosion mechanism of the as-cast, as-annealed, and hot-rolled NAB samples in 3.5 wt.% NaCl solution. All the samples consist of α matrix and κ phases (κ_III in as-cast NAB sample, whereas κ_III and κ_V in samples under the other two conditions), where the Al-rich κ phases are protected in a neutral 3.5 wt.% NaCl solution due to the formation of Al₂O₃ film on their surfaces [19]. The α phase becomes anodic to the κ phases and is preferentially corroded. The corrosion product film containing both aluminum and copper oxides is formed as a barrier to corrosive medium [59]. The corrosion product film continuously grows with the increase of immersion time and finally covers the surface completely, protecting better the NAB substrate [60]. The undissolved κ phases are retained in the corrosion product film [32].

The undissolved κ phases are detrimental to the protective property of the corrosion product film. On the one hand, the interfaces between the uncorroded κ phases and their surrounding corrosion products are susceptible, and the chloride ions can easily enter along with these interfaces [32], resulting in damage in the protective property of the corrosion product film and further corrosion of the substrate. On the other hand, the uncorroded κ phases also increase the growth stress of the corrosion product film and the risk of crack formation [32].

For the as-cast NAB sample, the large and continuous κ_III phases are retained in both the film and the matrix (Fig. 10(a)). For the as-annealed NAB sample (Fig. 10(b)), the interfaces between the κ phases and the surrounding corrosion products increase due to the growth of the continuous κ_III phases and the precipitation of a large number of κ_V phases, resulting in the severe damage of the film and the formation of more corrosion products. Therefore, the corrosion resistance of the as-annealed sample is worse than the as-cast sample. For the hot-rolled NAB sample (Fig. 10(c)), the κ phases are spheroidized and discontinuous in the α matrix. It is difficult for the chloride ions to further corrode the matrix along with the interfaces between the κ phases and the corrosion products, resulting in a decrease of the maximum local corrosion depth. In addition, hot rolling can eliminate defects, which is beneficial for the formation of a more stable passive film on the surface.

Besides, in the as-annealed sample, a large number of uncorroded κ phases (the continuous lamellar κ_III and lath-shaped κ_V) are densely distributed in the corrosion product film (Fig. 9(b)), leading to a great increase in the growth stress of the corrosion product film [32]. This induces the formation of cracks observed in the corrosion product film [61]. However, after further hot rolling, the original continuous κ phases (κ_III and κ_V) are spheroidized, decreased in content, and dispersed in the α matrix (Fig. 2(c)). When immersed in NaCl solution, these spheroidized κ phases are sparsely distributed in the corrosion product film, as shown in Fig. 9(c). The stress concentration caused by the spherical phases is less than that caused by lamellar or lath-shaped phases. Thus, the growth stress inside the film is smaller than that of the as-annealed sample.

The microstructural evolution, mechanical properties, and corrosion behavior of the as-cast, as-annealed, and hot-rolled NAB alloys (Cu-9Al-10Ni-4Fe-1.2 Mn) in 3.5 wt.% NaCl solution have been investigated. The main findings are drawn as follows:

(1) The microstructure of the as-cast NAB alloy contains α phase, κ_III phase, and some coarse NiAl strips. Annealing causes a large number of κ_V phases to precipitate in the α phase, and the size of the eutectoid κ_III increases after annealing. However, after further hot rolling, the original continuous κ phases (κ_III and κ_V) are spheroidized, dispersed, and decreased in the content; dislocations and twins are also formed.

(2) The as-cast NAB alloy exhibits the lowest strength and hardness, while the highest elongation. Annealing increases the strength and hardness while the decrease in elongation. However, further hot rolling simultaneously increases the strength, hardness, and elongation of the NAB alloy, where the yield strength is increased by ~84 % compared to the as-annealed alloy.

(3) Hot rolling improves the corrosion resistance of the NAB alloy. The values of mass loss increase from 3.53 ± 0.08 g/m² (as-cast) to 14.38 ± 0.31 g/m² for the as-annealed alloy, but decrease dramatically to 1.86 ± 0.09 g/m² for the hot-rolled alloy after immersion in the 3.5 wt.% NaCl solution for 28 days.

This work was supported by the National Key Research and Development Program of China (No. 2017YFA0403803), the National Natural Science Foundation of China (Nos. 51525401, 51834009, 51927801, 51901034, 51974058), and the Liaoning Revitalization Talents Program (No. XLYC1808005).