Fresh insights into structure–function-integrated self-antibacterial Cu-containing Al alloys: giving Al alloys a new function

Abstract

Contact infection by bacteria and viruses is a serious concern to human health. The increasing occurrence of public health problems has stimulated the urgent need for the development of antibacterial materials. Al alloys are the fastest-growing mass-produced material group, a prerequisite for the lightweight design of vehicles, food containers and storage, as well as civil-engineering structures. In this work, the structure–function-integrated concept was used to design and produce self-antibacterial Al–xCu (x = 2.8 and 5.7) alloys for the first time ever. The antibacterial tests indicated that Al–2.8Cu and Al–5.7Cu alloys provided a stable and efficient bacteriostatic rate against S. aureus and E. coli, which was 87% for Al–2.8Cu and 100% for Al–5.7Cu against S. aureus at 24 h, and 89% for Al–2.8Cu and 94% for Al–5.7Cu against E. coli at 24 h. The antibacterial effect was similar to the commonly-used antibacterial materials with a similar Cu content. Furthermore, the mechanical properties and corrosion resistance of Al–2.8Cu and Al–5.7Cu were comparable to those of the current commonly-used commercial casting Al–Cu alloys. Structural insights into the performance and biomedical function by Cu-rich precipitates provided understanding of the mechanisms of these structure–function-integrated self-antibacterial Cu-containing Al alloys: (i) the Cu-rich precipitates produced strengthening, and (ii) the immediate contact with Cu-rich precipitates and the Cu²⁺ caused a synergistic action in improving antibacterial activity. This work gives Al alloys a new function and inspires fresh insights into structure–function-integrated antibacterial Al alloys.

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New concepts

Contact infection by bacteria and viruses is a serious concern to human health. The increasing occurrence of public health problems has led to an urgent requirement to develop antibacterial materials. Al alloys are the fastest-growing mass-produced material group, a prerequisite for the lightweight design of vehicles, food containers and storage, as well as civil-engineering structures. In this work, a structure–function-integrated concept was used to design and produce self-antibacterial Al–xCu (x = 2.8 and 5.7) alloys for the first time ever. This work gives Al alloys a new function and inspires fresh insights into structure–function-integrated self-antibacterial Al alloys.

1. Introduction

Public health and safety issues are increasingly attracting attention. A key factor contributing to mass infection is the persistence of microbial pathogens on contact surfaces for an extended period of time.¹ In crowded environments like hospitals, pathogen adhesion and survival on touchable surfaces pose a serious threat to the population. In such situations, abiotic surfaces frequently touched by numerous individuals (such as door handles, handrails, elevator buttons, etc.) are important secondary hosts for cross-contamination and pathogen transmission.² The US center for disease control and prevention claimed that Escherichia coli (E. coli) causes the most frequent human bacterial infections and Staphylococcus aureus (S. aureus) is the second most frequent.³ Sessile microorganisms or pathogenic bacteria on abiotic surfaces have caused calamitous consequences and a significant cost in various fields such as public health place, food processing, health and medical-care occupations, home appliances, household supplies, etc.^4–8 The contact infection of bacteria is a serious concern to human health, which promotes an urgent requirement for developing antibacterial materials, especially antibacterial metallic materials, currently widely used bio-functional materials due to their characterized stable and reliable mechanical properties and good formability.^9–11

An effective approach to overcoming these challenges is to inhibit bacterial adherence and biofilm formation on the abiotic surface.¹² In this context, antibacterial metallic materials are highly favored because of the risk of bacterial infection. So far, various efforts have been explored to develop antibacterial metallic materials. Copper (Cu) has been used for medical applications due to its essential cofactor of several enzymes and good antibacterial and antifungal effects.^13,14 It is highly desirable to produce antibacterial metallic materials that can play an antibacterial role by alloying with antibacterial Cu. The release of Cu²⁺ from Cu-containing alloys may produce the desired antibacterial effect. The antibacterial mechanism of Cu²⁺ was demonstrated to be attributed to oxidative stress, membrane damage, protein denaturation, etc.¹⁵ Several types of Cu-containing stainless steels, Cu-bearing high-entropy alloys, Fe–Cu alloys, Mg–Cu alloys, and Ti–Cu alloys were developed, which showed reliable antibacterial activity against common pathogens.^9,16 However, these Cu-containing alloys have disadvantages related to their physical, chemical and mechanical behavior, their production processing, and their economic cost for these potential commercial applications.¹⁷

In contrast, Al alloys, another popular structural material, belong to the fastest-growing mass-produced metal group, a prerequisite for lightweight design of vehicles, food containers and storage, as well as civil-engineering structures.^18–20 Applying a protective layer to a surface devoid of antibacterial activity is the preponderant antibacterial modification technique employed today for Al alloys.²¹ Zhang et al.²² initially generated an anodic aluminum oxide (AAO) using an oxalic acid solution, then electrodeposited Cu using a copper sulfate solution to produce an antibacterial layer on the surface of Al alloys. Cerchier et al.²³ created a high-quality, evenly-distributed antibacterial silver plasma electrolytic oxide (PEO) coating on the surface of Al alloys, utilizing an alkaline solution of silicate compounds and silver micron particles to achieve antibacterial activity, and also to improve corrosion resistance. Existing research has mainly focused on antibacterial coating by surface modification technology to achieve Al alloys with antibacterial activity. Although current technology could prepare Al alloys with coatings with good antibacterial activity, the coatings are brittle with weak binding to the Al alloys, with the consequence that the coating is easy to break and falls off during use.²⁴ Furthermore, coating is difficult for Al alloys of complex shape, resulting in poor uniformity of the coating.²⁵ In these cases, the antibacterial problem of the Al alloy was not effectively solved. In addition, the antibacterial effectiveness of such surface-treated Al alloys was not stable, and decreased as the antibacterial film decayed. Pornnumpa et al.²⁶ used an in situ chemical reduction process to deposit silver nitrate on the surface of anodized Al alloys. The antibacterial rate against E. coli was reduced from 100% to just 72% on the 90th day. Evidently, it is challenging to achieve a long-term sustainable antibacterial activity for surface-treated Al alloys. The previous antibacterial research on producing antibacterial metals by Cu alloying inspired the aim of producing self-antibacterial Cu-containing Al alloys by alloying Cu into Al, in which Cu provides the antibacterial activity. This approach can be expected to achieve the desired long-term sustainable antibacterial activity for widely-used Al alloys devoid of antibacterial activity. However, to the best of our knowledge, never before has any work been conducted to develop such self-antibacterial Cu-containing Al alloys, and furthermore, the details regarding their antibacterial mechanism are not clear. In this type of self-antibacterial Cu-containing Al alloys, the content and distribution of Cu-rich intermetallic compounds (IMCs) are expected to be the critical factors affecting their antibacterial activity, and mechanical and corrosion properties.

In this work, a structure–function-integrated concept was used to design and produce self-antibacterial Al–xCu (x = 2.8 and 5.7) alloys for the first time ever. The mechanical, corrosion, and antibacterial properties of these alloys were evaluated to appraise their prospects for application in public health places, food processing, health and medical-care and household supplies, etc., paying special attention to the influence of the content and distribution of Cu-rich IMCs on antibacterial activity. This work gives Al alloys an unprecedented new function and extends their applications in fields with antibacterial requirements.

2. Materials and methods

2.1. Materials

The experimental Al alloys were designed with the nominal compositions of 2.8 wt% Cu (hereafter, Cu2.8) and 5.7 wt% Cu (hereafter, Cu5.7). The Cu free alloy (hereafter, Cu0) was used as a blank control for antibacterial experiments. Cu5.7 was chosen because the ultimate solid solubility of Cu in Al is 5.65 wt% according to the Al–Cu binary phase diagram, and Cu2.8 is half this value. Table 1 lists the chemical composition of the Al alloys, evaluated using an inductively coupled plasma optical emission spectrometer (ICP-OES). Cu2.8 and Cu5.7 were ZL201 Al alloys, a type of Al alloy with a high Cu concentration, while Cu0 with consistency in other components is a de-coppered Al alloy that served as a blank control. The ingots were melted in a well type resistance furnace at 760 °C. An appropriate amount of covering agent (NaCl : KCl : Na₃AlF₆ = 1 : 1 : 2) was added after the furnace was opened every time to isolate the air. After melting, the melt was stirred to ensure a homogeneous distribution of alloying elements, and poured into an iron mold that had been coated with ZnO and dried at 200 °C, and air cooled to room temperature. The Cu2.8 and Cu5.7 specimens cut from the cast ingots were subjected to T6 heat treatment: solid solution (at 525 °C for 16 h followed by water quenching) plus aging (at 170 °C for various times to reach peak aging followed by water quenching). The pertinent statistics indicated that the peak aging times were 8 h and 14 h for Cu2.8 and Cu5.7, respectively.

Table 1 Chemical compositions of the experimental Al alloys evaluated using ICP-OES (wt%)

Samples no.	Cu	Mn	Ti	Fe	Al
Cu0	0	0.69	0.35	0.10	Balance
Cu2.8	3.00	0.67	0.33	0.10	Balance
Cu5.7	5.71	0.65	0.35	0.09	Balance

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The microstructure was characterized using a scanning electron microscope (SEM) with an energy spectrometer (EDS) and a transmission electron microscope (TEM). SEM observations were carried out using a TESCAN MIRA4 LMH instrument equipped with One Max 50 EDS. TEM observations were carried out using an atomic resolution Thermo Fisher Scientific Talos^TM F200X STEM operated at 200 KV. Specimens for SEM were ground using silicon carbide abrasive papers and mechanically polished using diamond paste. Specimens for the TEM observations were prepared by mechanically grinding to a final thickness of 60–80 μm, and 3 mm thin foil discs were punched out for twin-jet electro-polishing in a solution of 25 vol% nitric acid in methanol at 18 V and at −25 to −30 °C. X-ray diffraction (XRD) was performed to identify the phase compositions at a scan step of 4° min⁻¹ from the range of 20° to 80° using an X-ray diffractometer (XRD: D/Max 2500).

The room temperature tensile experiments were conducted at a constant speed of 2 mm min⁻¹ according to the ASTM E8M standard. The gage width was 8 mm and the length of the parallel section was 40 mm. Three specimens of each category were tested to obtain an average value. The fracture morphologies of the specimens after the tensile test were examined using the SEM.

2.2. Corrosion evaluations

Corrosion behavior of the Al alloys was evaluated using electrochemical tests and the neutral salt spray test. Specimens had dimensions for electrochemical corrosion evaluations of 10 mm × 10 mm × 2 mm and for a neutral salt spray test of 15 mm × 10 mm × 2 mm, were mechanically ground to 2000-grit with SiC papers and were mechanically polished. Three specimens of each alloy were examined.

Potentio-dynamic polarization curves and electrochemical impedance spectroscopy (EIS) were carried out using the MULTIAUTOLABM204 electrochemical workstation with a 3.5% NaCl aqueous solution, which was made with analytical grade reagent and distilled water. A platinum gauze was used as the counter electrode. A saturated calomel electrode (SCE) was used as the reference electrode. The polished specimens were the working electrode, which were encapsulated into epoxy resin with an exposed surface area of 10 mm × 10 mm. All potentials were referred to the SCE. The scanning rate was 0.1 mV s⁻¹ from −1 to −0.1 V after the specimens were held at the open circuit potential for 30 min to reach the steady state. The EIS was performed at the open circuit potential with a frequency range of 10⁻²–10⁵ Hz.

The neutral salt spray test was performed using the DIN EN ISO 9227 salt spray corrosion testing machine according to the GB/10125-1997 standard. The temperature was 35 °C, the pH was 6.8, the NaCl solution concentration was 5%, and the experimental period was 168 h. There were three parallel specimens in each group. The specimens were placed on a V-shaped frame with the test surface facing upwards and at an angle of 30° to the horizontal direction. On the 0th, 1st, 3rd, 5th, and 7th day in the experiment period, the specimens were removed, dried indoors for 30 min, and observed using a camera and an optical microscope to characterize the corrosion status of the surface. The corrosion morphology of specimens after 7 days of salt spray was observed using SEM.

2.3. Antibacterial assay

The antibacterial properties were assessed using E. coli and S. aureus plate counts. The sterilized samples (Cu0, Cu2.8, and Cu5.7) were placed in 24-well culture plates. Bacterial incubation was conducted on BKMAM solid agar medium, with a bacterial suspension prepared at a concentration of 10⁵ cfu per L. A 100 μL aliquot of this suspension was uniformly distributed on the surface of each sample. The 24-well plates were incubated in a controlled thermostat (FYL-YS-280L) for predetermined intervals (12, 24, 36, and 48 h). The bacterial suspension on the sample surface was pipetted out and applied to the plate that was inverted and incubated in a thermostat. After the incubation, photographs were taken. The antibacterial rate was calculated:

K = (A − B)/A × 100% (1)

where K is the antibacterial rate, A is the number of bacterial colonies in the blank control group, and B is the number of bacterial colonies in the experimental group. Three parallel experiments were conducted for each group.

The samples were co-cultured with S. aureus for 24 h and then fixed in 2.5 vol% glutaraldehyde for 4 h. The samples were dehydrated gradually by immersion in increasing concentrations of alcohol (50, 60, 70, 80, 90, 95 and 100 vol%). The sterile PBS was rinsed 5 times, and dried at room temperature. The samples were stained using the LIVE/DEAD BacLightTM Bacterial Viability Kit (FUSHENBIO) under light-protected conditions. Live bacteria were stained green by SYTO-9 and dead bacteria were stained red by propidium iodide (PI). The biofilm thickness was evaluated by creating a digital 3D profile of the bacterial biofilm using a confocal laser scanning microscope (CLSM: Zeiss AirScan). SEM (TESCAN MIRA3) was used to characterize the bacterial morphology on the sample surface.

To study the release of the Cu²⁺, the specimens were immersed in deionized water with a surface area to volume ratio of 0.8 cm² mL⁻¹ at room temperature according to the GBT 16886.12-2017 standard. The content of Cu²⁺ was measured using inductively coupled plasma mass spectrometry (ICP) after the specimens were immersed for 12 h and 24 h, respectively.

3. Results

3.1. Microstructures

Fig. 1 shows the SEM-EDS results of Cu2.8 and Cu5.7. Every microstructure was mainly composed of the α (Al) matrix and some intermetallic compounds (IMCs). Typical regions of IMCs as marked (Fig. 1a and b) were analyzed using EDS. The EDS results are semi-quantitative to indicate the compositions. Point A in Fig. 1a₁ and Point B in Fig. 1b₁ are highlighted as representatives. The corresponding compositions evaluated using EDS were: Point A mainly contained Al, ∼21 wt% Cu and ∼4 wt% Mn; Point B mainly contained Al, ∼21 wt% Cu and ∼8 wt% Mn. The EDS maps of Al, Cu and Mn in the marked regions revealed that the coarse IMCs were (Cu, Mn)-rich, as shown in Fig. 1a₃–a₅ and b₃–b₅. Therefore, the IMCs in Point A and Point B were composed of the (Al, Cu, Mn) phase. These coarse (Al, Cu, Mn)-containing IMCs were non-equilibrium eutectic insoluble phases formed during solidification.^27,28 In addition, there were some sub-micron and nano sized IMCs in the α (Al) matrix.

Fig. 1 SEM-EDS data: ((a) and (a₁)–(a₅)) for Cu2.8, and ((b) and (b₁)–(b₅)) for Cu5.7.

Fig. 2a and b show the TEM micrographs of Cu2.8 and Cu5.7, which were taken near the 〈001〉_Al direction. There are numerous IMCs in the α (Al) matrix. The IMCs are mostly rodlike, have specific orientations and regular shapes with diameters of 20–100 nm and lengths of 500–1000 nm. The subcircular IMCs are a cross-section of the rodlike IMCs. The axis direction of these IMCs is schematically shown in Fig. 2c, consistent with the 〈001〉_Al zone axis. EDS point analysis (Fig. 2a₁ and b₁), representing sites in typical IMCs, showed that Point A₁ mainly contained Al, ∼7.4 at% Cu and ∼11.1 at% Mn; Point B₁ mainly contained Al, ∼7.9 at% Cu and ∼13.5 at% Mn, which indicated an atomic ratio of Cu : Mn ∼2 : 3. Selected area electron diffraction (SAED) further confirmed that the phase was a T (Al₂₀Cu₂Mn₃) phase (Fig. 2a₂ and b₂). Cu2.8 was used as a representative to study the element distribution. EDS maps revealed that the IMCs were rich in Cu and Mn, as shown in Fig. 2d₁–d₄. Therefore, the IMCs were composed of Al, Cu, and Mn.

Fig. 2 TEM, EDS and SAED: (a) Cu2.8, (b) Cu5.7, (c) sketch of IMC distribution in the Al matrix, and (d₁)–(d₄) mapping result for Cu2.8 (Al is red, Cu is green, and Mn is yellow).

The structure of the rod-shaped T-phases was investigated using a high resolution transmission electron microscope (HRTEM). Fig. 3a and 4a depict the interfaces between the matrix and typical rod-shaped T-phases in Cu2.8 and Cu5.7, respectively. Inverse fast Fourier transform (IFFT) images of the matrix and T phases in Cu2.8 (as shown in Fig. 3b₁, b₂ and c₁–c₃) revealed that the (020) and (101) planes of the T phases were parallel to the (020) and (002) planes of the Al matrix, respectively (defined as orientation relationship I: OR-I), i.e. (020)_T//(020)_Al and (101)_T//(002)_Al (as shown in Fig. 3e). The IFFT image of the interface (as shown in Fig. 3d) indicated that the (060)_T of the T-phase was epitaxial to the (020)_Al of the matrix. However, in Cu5.7, the orientation relationship of the T-phases and matrix (OR-II) differed somewhat, with (020)_T//(020)_Al and (200)_T//(200)_Al (as illustrated in Fig. 4e). The T-phase had an orthorhombic lattice, with a = 2.41 nm, b = 1.25 nm, and c = 0.74 nm. The angle between (101)_T and (002)_T was 17.22°. Therefore, the difference between OR-I and OR-II was that the crystal cell of Al rotated 17.22° about the [010]_T axis. The T-phase exhibited a B-centered orthorhombic structure.^29,30 The corresponding cell structure is shown in Fig. 4f, which was consistent with the literature.^29,30

Fig. 3 (a) HRTEM image of the interface between the T phase and the matrix in Cu2.8, IFFT images of the matrix (b₁) and (b₂), T-phase (c₁)–(c₃), and interface (d), and (e) the sketch map of the orientation relationship between the T-phase and Al.

Fig. 4 (a) HRTEM image of the interface between the T phase and the matrix in Cu5.7, IFFT images of the matrix (b₁) and (b₂), T-phase (c₁)–(c₃), and interface (d), (e) sketch map of the orientation relationship between the T-phase and Al, and (f) a schematic diagram of the T-phase structure.

In addition, the difference in Cu content also led to differences in the precipitation of θ phase particles in the matrix. TEM images of a T-phase in Cu2.8 and Cu5.7 are displayed in Fig. 5a and b. There were moiré fringes along the [010]_T direction in the T-phase, indicating that there was overlap between the T-phase and the matrix along this direction, and the lattice constants of the two phases were close,³¹ as depicted in Fig. 3 and 4. Although the rod-shaped T-phases had comparable geometries, the Al matrix varied significantly. Cu5.7 contained tiny evenly-distributed dispersions. These IMCs had a length of around 10–30 nm, a width of less than 2 nm, were needle-shaped and were orthogonally dispersed in the matrix (as shown in Fig. 5c). High-angle annular dark field (HAADF) images of the needle-like IMCs are presented in Fig. 5d. The structure had three layers of Al atomic planes separating every two layers of Cu atomic planes, indicating that these needle-like IMCs were the θ′′ phase.³² However, the Al matrix of Cu2.8 alloy within the field of view in Fig. 5a was smooth, rather than a mesh matrix of the Cu5.7 alloy. We attempted to enlarge the Al matrix, but did not find any θ′′ phases in the Cu2.8 alloy. Fig. 5e shows the cross-section of a T-phase particle in Cu5.7. This typical sample indicated that the cross-sectional structure displayed a twinning structure. An enlarged image of the twin zone revealed that the T-phases were produced by a parallel inlay of flat hexagonal substructure units, whereas the twin structure consisted of intersecting flat hexagons. Due to the presence of twinning, the interface relationship between the T phase and the matrix was complex. Although the orientation relationship between each twin and the matrix was different, the interface between them and the matrix was the same, i.e. {200}T and {101}T, as shown in Fig. 5e. In addition, near the cross-section, there were needle-like phases with a width of ∼10 nm. The magnified image of these needle-like phases, as shown in Fig. 5f, represents the atomic configuration of the θ′ phases.³³ The SAED image also confirmed that they were θ′ phases.³³ The complete precipitation sequence of the θ series phases was: supersaturated solid solution → GP zone → θ′′ phase → θ′ phase → θ phase.^33,34 On the (002)_Al surface, a layer of Cu atomic plane initially formed in the early stages of aging and was entirely coherent with the Al matrix. There was some shrinking close to the Cu atomic plane as a result of the difference in atomic radii between Cu and Al. The growing GP area caused the θ′′ phase (Al₃Cu), which exhibited perfect coherence with the matrix. The structure of the θ′′ phase is shown in Fig. 6a, where Cu atoms occupy the (001)_Al atomic plane of the matrix, with three levels of Al atomic planes separating every two Cu atomic plane layers. The generation of the semi-coherent θ(Al₂Cu) phase occurred as the aging process progressed, as shown in Fig. 6b. Ultimately, it changed into a stable θ(Al₂Cu) phase with a complicated orientation relationship and a crystal structure that was non-coherent with the matrix, as shown in Fig. 6c. In Cu5.7, solute atoms were desolvated and precipitated from the supersaturated solid solution in the quenched state to form atomic clusters (GP zone), and because of the high content of Cu atoms near the T phase, the GP zone was preferentially nucleated near the T phase. Therefore, at the peak aging, it grew into the θ′ phase near the T phase, and the uniformly distributed precipitated phase in the matrix was the θ′′ phase. However, the Cu concentration in Cu2.8 was low, and a substantial quantity of T phase was generated during solid solution, leading to depletion of Cu in the matrix, resulting in θ′′ precipitation inhibition during aging.

Fig. 5 TEM images of high magnification: (a) Cu2.8, (b) Cu5.7, (c) needle shaped θ′′ phases and HAADF image (d) in Cu5.7, (e) T-phase cross-sectional, and (f) HAADF image and SAED of the θ′ phase.

Fig. 6 Precipitation sequence of the theta-series phases: (a) θ′′ phase, (b) θ′ phase, and (c) θ phase.

3.2. Mechanical properties

The engineering stress–strain curves and associated mechanical characteristics at room temperature are presented in Fig. 7. The Cu content has a significant impact on the mechanical characteristics as an essential element of IMCs in the Al–Cu alloys. Cu5.7 had high strength, with a yield strength (YS) 102% higher and an ultimate tensile strength (UTS) 66% higher compared to Cu2.8. As strength increases, plasticity decreases. The elongation of the Cu2.8 (14%) was larger than that of the Cu5.7 (7%). The homogenous dispersed precipitation of θ′′ in the matrix was the primary cause for the increased strength of Cu5.7 compared to Cu2.8.

Fig. 7 (a) Stress–strain curves and (b) mechanical properties.

The fracture morphologies of Cu2.8 and Cu5.7 are presented in Fig. 8a and b. The room-temperature tensile tests showed clear ductile fracture. The fracture was mainly composed of dimples of varying sizes and tear ridges and a few cleavage planes. Cu5.7 had small and shallow dimples at fracture, and was less plastic and had lower elongation, while the Cu2.8 had large and deep dimples at fracture, suggesting higher plasticity.³⁵

Fig. 8 Tensile fracture surface morphology of Cu2.8 (a) and Cu5.7 (b).

3.3. Electrochemical properties and neutral salt spray tests

Fig. 9a shows the polarization curves after the open circuit potential was stabilized in the 3.5% NaCl solution. The polarization curves were fitted using the Tafel extrapolation method. The results are shown in Table 2. In general, a higher corrosion current density indicates a faster corrosion rate.^36–38 Cu2.8 had a lower corrosion current density of 8.6 μA cm⁻², indicating a lower corrosion rate, and a higher corrosion resistance.

Fig. 9 (a) Polarization curve of Cu2.8 and Cu5.7 in 3.5% NaCl solution, (b) Nyquist diagram and partial enlarged perspective, (c) Bode diagram, and (d) equivalent circuit fitting based on EIS.

Table 2 Corrosion potential E_corr and corrosion current density i_corr of Cu2.8 and Cu5.7 in 3.5 wt% NaCl solution based on the polarization curves

Samples no.	E corr (V)	i corr (μA cm^−2)
Cu2.8	−0.60	8.62
Cu5.7	−0.58	10.16

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Electrochemical impedance spectroscopy (EIS) was performed in 3.5% NaCl solution to evaluate the electrochemical corrosion performance of the Al alloys. Fig. 9b and c present the Nyquist and Bode curves. In general, the diameter of the capacitive arc relates to the corrosion resistance, with a larger diameter indicating larger corrosion resistance. The capacitance arc diameter of the EIS curve of Cu2.8 was larger, indicating higher corrosion resistance.^39,40 There were two arcs in the Nyquist curve of each alloy. The capacitive arc in the medium-high frequency range was associated with the formation of an oxide film on the specimen surface. The capacitive arc in the medium-low frequency range was associated with the electrochemical response of the local corrosion process.⁴¹ The Bode curves of all specimens had two peaks, one in the mid-high frequency region with a phase angle of about 70–80° and the other in the mid-low frequency band with a phase angle of around 30–40°. These two peaks correspond to the two capacitive arcs in the Nyquist curve and had the same significance. Cu5.7 had a smaller impedance modulus (|Z|) in the low-frequency range, suggesting lower corrosion resistance,^40,42 which followed the same pattern as the Nyquist plot.

An equivalent circuit (EEC) was utilized to fit the EIS. The EEC model is presented in Fig. 9d. The fitting was done with ZSimpWin. R_s is the electrolyte solution resistance; Q_f and R_f are the capacitance and resistance of the oxide film produced on the specimen surface, respectively; Q_ct and R_ct are the interfacial double layer capacitance and charge transfer resistance of the specimen surface (inner layer of the oxide film), respectively. Q_ct and Q_f are constant phase angle elements describing the non-ideal capacitive behavior.⁴³ The results of the fit are shown as solid lines in Fig. 9b and c.

Table 3 shows the fitted parameters for the EIS data. The polarization resistance R_p (R_p = R_f + R_ct) may be used to calculate the corrosion resistance.⁴⁴ Increasing Cu content decreased R_p from 38.9 to 14.7 KΩ cm⁻², increased Q_f from 8.1 to 39 μF cm⁻², and increased Q_ct from 80.3 to 255.4 μF cm⁻². All these indicated that increasing Cu content reduced the protective ability and corrosion resistance of aluminum alloy passivation films.^45,46

Table 3 Equivalent circuit parameters of the two alloys fitted to the EIS

Samples no.	R (Ω cm^−2)	Q f (μF cm^−2)	n f	R f (kΩ cm^−2)	Q ct (μF cm^−2)	n ct	R ct (kΩ cm^−2)
Cu2.8	14.59	8.10	0.89	7.95	80.30	0.79	30.98
Cu5.7	6.88	39.00	0.92	4.81	255.4	0.98	9.85

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Fig. 10 and 11 depict the surface macroscopic morphologies following 7 days of salt spray using digital and optical images, respectively. As shown in digital photos (Fig. 10), the images of the alloys both had a normal gray luster. As shown in optical images (Fig. 11), the corrosion active spots occasionally appeared and the pitting corrosion pits sporadically distributed on the surface of both alloys. With the increase in salt spray test time, the number of pitting corrosion pits on the surface of samples increased while keeping the size that stopped developing. The Al alloy matrix in the small area around the pitting corrosion pits remained still free of corrosion, consistent with the characteristic morphologies of Al alloy corrosion.^47,48 This indicated that the corrosion of the Al alloy was anodic dissolution type, in which the area around the pitting was cathodic protection, and the inside of the pitting was anodic corrosion. It is well known that the surface of Al alloy can easily induce pitting corrosion in the presence of chloride ion,^19,47,48 and hence pitting corrosion of Al alloys is unavoidable in this case. Although the pitting corrosion happened for both alloys in the present salt spray test, the size of the small pitting pits did not develop continuously when increasing the salt spray test time, and the non-pitting surface remained still free of corrosion, indicating that the present alloys have a certain corrosion resistance in the salt spray corrosion test environments. Furthermore, the surface pitting corrosion macroscopic morphologies indicated the Cu2.8 alloy had more corrosion resistance than the Cu5.7 alloy.

Fig. 10 Digital images of the salt spray test of the Cu2.8 and Cu5.7 after 7 days.

Fig. 11 Optical images of the salt spray test: (a₁) 1 day, (a₂) 3 days, and (a₃) 7days for Cu2.8; (b₁) 1 day, (b₂) 3 days, and (b₃) 7days for Cu5.7. Red arrows in the images marking representative pitting corrosion pits.

Fig. 12 shows the SEM images of the corrosion morphology on the alloy surface after 7 days of salt spray. The low magnification morphology images in Fig. 12a and c indicate that the surfaces of both alloys were subjected to varying degrees of corrosion. The corrosion of Cu5.7 was more severe than that of Cu2.8, consistent with the electrochemical data. There were corrosion products and corrosion rings. The high magnification images in Fig. 12b and 10d indicate that the intermetallic compounds originally distributed along the grain boundaries had peeled off. Fig. 12e₁–e₅ show an EDS analysis of a corrosion pit of Cu5.7, with a diameter of ∼250 μm. There was substrate detachment at the edge. The distribution of O was consistent with the enrichment of Cu and Mn, indicating that corrosion was more likely to occur near the Cu and Mn phases.

Fig. 12 Corrosion morphology after 7 days of neutral salt spray: (a) and (b) Cu2.8, (c) and (d) Cu5.7, and (e₁)–(e₅) EDS analysis of a corrosion pit of Cu5.7.

The performances of these alloys were compared to the normally used Al–Cu alloys,^47–51 with tensile strength and corrosion current density serving as measuring markers for mechanical qualities and corrosion resistance. Fig. 13 indicates that Cu2.8 and Cu5.7 fulfill the requirements of mechanical performance and corrosion resistance as structural materials.

Fig. 13 Comparison between these alloy and common Al–Cu alloys.

3.4. Antibacterial properties

Fig. 14 shows the plate colony count images of E. coli and S. aureus co-cultured with Cu0, Cu2.8, and Cu5.7 at room temperature for 12 h, 24 h, 36 h, and 48 h, with plates that do not come into contact with any alloy as blank controls to calculate the antibacterial rate. The results are shown in Fig. 15. The antibacterial tests indicated that Cu2.8 and Cu5.7 provided a stable and efficient bacteriostatic rate against S. aureus and E. coli, which was 89% for Cu2.8 and 94% for Cu5.7 against E. coli at 24 h (Fig. 15a), and 87% for Cu2.8 and 100% for Cu5.7 against S. aureus at 24 h (Fig. 15b).

Fig. 14 Images of plate colony counts of E. coli and S. aureus after 12 h, 24 h, 36 h, and 48 h of co-cultivation with Cu0, Cu2.8, and Cu5.7.

Fig. 15 Antibacterial rate after 12 h, 24 h, 36 h, and 48 h co-culture of Cu0, Cu2.8, and Cu5.7 with (a) S. aureus and (b) E. coli.

The morphologies of the colonies for Cu0 indicate that there were numerous colonies and that they were dispersed uniformly throughout the plate. The surface of Cu0 supported the growth of microorganisms, suggesting that Cu0 lacked antibacterial qualities. The culture media of these samples indicated a noteworthy decrease in the number of bacteria for the surface of Cu2.8 and Cu5.7 as compared to the experimental reference Cu0. This suggested that the addition of Cu element efficiently eliminated bacteria. The fewer bacterial colonies on Cu5.7 than on Cu2.8 at any testing stage suggested that the antibacterial activity of the Cu-containing Al alloys was dependent on the amount of Cu.

Wang et al.⁵² discovered that after annealing at 740 °C, the antibacterial rates of Ti–3Cu and Ti–5Cu against S. aureus were 97% and 99%, respectively. Xi et al.⁵³ discovered that the antibacterial rate of 316L–2.5Cu stainless steel cocultured with S. aureus for 24 h was 95%. Ji et al.¹⁵ found that selective laser melted (SLM) Ti–3Cu had a 99% antibacterial rate against E. coli after 24 h. The antibacterial capabilities of the alloys in this study indicated that they had an antibacterial effect similar to Ti alloy or stainless steel with the same Cu concentration.

Fig. 16 depicts the morphology of E. coli and S. aureus cultured on the samples for 24 h using SEM, and distinguishes between dead and live bacteria using false colors. Red represents inactive bacteria, while yellow and green represent live bacteria that can grow normally. There were a large number of bacteria with good growth conditions on the surface of Cu0. The bacterial particles were full and clustered in large numbers, indicating that Cu0 did not have an antibacterial effect. Co-culturing the Al–Cu alloys with bacteria for 24 h produced fewer bacteria, and the bacterial shape was irregular, shriveled, and some bacteria had content flowing out. The entire cell structure had been destroyed. SEM images indicate that these Al–Cu alloys hindered the development and metabolism of S. aureus and E. coli, and their antibacterial properties were related to their copper content.

Fig. 16 Surface morphology of the alloys after co-culturing with S. aureus (a₁)–(a₃) and E. coli (b₁)–(b₃) for 24 h: (a₁) and (b₁) Cu0, (a₂) and (b₂) Cu2.8, and (a₃) and (b₃) Cu5.7.

The bacterial activity was evaluated using live/dead fluorescence staining after 24 h for these Al–Cu alloys and S. aureus co-cultivation. The copper-free Cu0 was the control group. The findings are illustrated in Fig. 17a–c. The surface of the control Cu0 co-cultured with S. aureus exhibited a high level of green fluorescence, while the red fluorescence on the sample surface was low, indicating that bacterial activity on the surface was strong, with almost no bacterial death, and the biofilm was relatively thick. In contrast, the surface of Cu2.8 exhibited an increase in the proportion of red fluorescence, indicating that some bacteria had lost viability. The surface of Cu5.7 emitted red fluorescence, whereas green fluorescence was reduced. The properties of red and green fluorescent dyes indicated that these copper-containing Al alloys had good antibacterial effects against S. aureus after 24 h of cultivation, and that the antibacterial activity was positively related to the Cu content.

Fig. 17 CLSM images of bacterial live/dead fluorescent staining (a)–(c) and release of Cu2⁺ (d).

The antibacterial mechanism of these Al–Cu alloys was studied by the measurement of the release of Cu²⁺ from Al–xCu (x = 0, 2.8, and 5.7) in deionized water using ICP. The specimens were immersed in deionized water with a surface area to volume ratio of 0.8 cm² mL⁻¹ at room temperature. The content of Cu²⁺ was measured using ICP after the specimens were immersed for 12 h and 24 h, respectively. Fig. 17d shows the release of Cu²⁺ at different soaking times (12 h and 24 h). After soaking in deionized water, Cu²⁺ was detected in the Al alloys containing Cu, and the amount of Cu²⁺ was increased with the Cu content. When soaked for 24 h, the amount of Cu²⁺ released by Cu5.7 was 20 μg L⁻¹, while the amount of Cu²⁺ released by Cu2.8 was 37 μg L⁻¹. There was a significant difference in the amount of Cu²⁺ between Cu2.8 and Cu5.7. As the soaking time increased, the amount of Cu²⁺ increased. However, the release of Cu²⁺ at 12 h was higher than half of the quantity at 24 h, indicating that the release of Cu²⁺ in deionized water was not continuous and uniform. In the early stages of soaking, the release rate of Cu²⁺ in the alloy was faster. Increasing soaking time decreased the release rate, due to the decrease in concentration difference and the surface passivation film hindering the dissolution of Cu²⁺.

4. Discussion

4.1. Model analysis of strengthening mechanisms

The strength of alloys depends on grain size, dislocation density, solid solution strengthening, and precipitated phase size and density.⁵⁴ These alloys were aged after casting, so their dislocation density was low. Taking into account potential strengthening mechanisms, the strength of these alloys may be computed using formula (2):⁵⁵

σ_YS = σ₀ + σ_GB + σ_ss + σ_ps (2)

where σ_YS is the total yield strength, and σ₀ is the yield strength of pure Al, with a value of 10 MPa. σ_GB, σ_ss and σ_ps are the increments in yield strength caused by grain boundaries, solid solution atoms, and precipitate precipitation, respectively.

4.1.1. Grain boundary strengthening. The increase in grain size decreases the yield strength, as indicated by the Hall–Petch equation:⁵⁶

σ_GB = kd^−1/2 (3)

where k is the Hall–Petch constant, 0.08 MPa m^1/2 for Al–Cu alloys, and d is the average grain diameter. The average grain sizes of Cu2.8 and Cu5.7 were 135 μm and 133 μm, respectively. The grain strengthening for Cu2.8 and Cu5.7 was 6.88 MPa and 6.94 MPa, respectively, according to eqn (3).

4.1.2. Solid solution strengthening. The solid solution strengthening is characterized quantitatively:⁵⁷where c_i is the mass fraction of the solid solution element in the matrix, k_i is the solid solution strengthening efficiency of the corresponding element, m is a constant, and m is usually assumed to be 1 in the Al–Cu–Mn alloy system. The solid solution strengthening in these Al–Cu–Mn alloys is mostly caused by Cu and Mn atoms dissolved in the matrix. The literature indicates that k_Mn = 30.3 MPa per wt% and k_Cu = 13.8 MPa per wt%.⁵⁸Table 4 displays the mass fractions of dissolved elements as determined by EDS spectroscopy in the α-Al matrix. Eqn (4) indicates that the contribution of Cu2.8 and Cu5.7 solid solution to the yield strength was 34 MPa and 53 MPa, respectively.

Table 4 Mass fraction of elements in solid solution in the Al matrix

Samples no.	Cu (wt%)	Mn (wt%)
Cu2.8	1.48	0.44
Cu5.7	2.35	0.49

---

4.1.3. Precipitation strengthening. The cutting or bypassing mechanisms for various precipitate shapes and sizes are determined by eqn (5)–(7), respectively, which are used to calculate their contributions to the yield strength:^59–61where M represents the Taylor factor of Al, G represents the shear modulus of Al, b represents the Burgers vector of Al, and ϑ represents the Poisson's ratio of Al, which are 3.06, 27.8 GPa, 0.286 nm, and 0.345, respectively. Δμ (fitting parameter, 4.03 GPa) represents the shear modulus between the phase and the substrate. The geometric constant A is related to the precipitation habit plane and is taken as 1.225 for the (100) plate. C is a constant, taken as 1.23, for a random array of precipitates. The internal cut-off radius for dislocation energy calculation r₀ is equivalent to 2b. λ is the average distance between the precipitates, and t_p represents the average diameter of precipitates. f_v is the volume fraction of precipitates, calculated using the following formula:^62,63N_v, N_p, D_m and D_p represent the number density of precipitates, the number of precipitates counted in the selected area, the average diameter (particle length), and the true diameter (related to D_m), respectively. t and A_s are used to represent the thickness of the Al foil and the area of the TEM images.

The TEM images of Fig. 2 and 5 show that a total of three forms of Cu-rich IMCs were involved in these alloys: fine θ′′ phases, rod-shaped T phases, and θ′ phases in the vicinity of the T phases. According to ref. 59, the small dispersions in Cu5.7 with a size of 20 nm (θ′′ phases) are the weaker precipitates, which can be cut by dislocations and are described by the strengthening model of eqn (5). The distribution of the precipitated phases in these alloys was statistically determined using ImageJ software. The volume fraction of the θ′′ phases in the Cu5.7 was 1.4%, which contributed 80 MPa to the yield strength. The considerable size of the T-phases indicates dislocation movement was hindered through the Orowan mechanism.⁶⁰ This effect can be estimated using eqn (6). The statistical result of IMCs size is shown in Table 5. The contribution of the IMCs to yield strength of these alloys was calculated to be σ_ps,_Cu2.8 = 66 MPa, and σ_ps,_Cu5.7 = 153 MPa.

Table 5 Size parameters of T phases and θ′′ phases

Sample no.	Precipitate	Mean length Dm (nm)	Mean thickness tp (nm)	Mean thickness f (%)
Cu2.8	T	600	90	3.43
Cu5.7	θ′′	18.7	1.45	1.40
	T	650	110	4.64

---

The contribution of each mechanism to yield strength for these alloys determined through calculation is shown in Fig. 18. σ_0.2 represents the yield strength obtained from the experiment. Precipitation strengthening was the main strengthening mechanism. The reason why the yield strength was somewhat lower than the calculated value was attributed to casting defects.

Fig. 18 Contribution of different strengthening mechanisms to the yield strength of these alloys.

4.2. Localized corrosion of alloys

Generally speaking, a dense oxide film on the surface of the Al–Cu alloys can effectively inhibit corrosion and improve the corrosion resistance.⁶⁴ However, in neutral salt spray environments, higher Cl⁻ concentrations and higher relative humidity can damage the surface oxide film, leading to corrosion.⁶⁵ The alloying content in these Cu containing Al alloys was relatively high, leading to a large quantity of IMCs. These IMCs were not consistent with the substrate potential and can easily form micro corrosion batteries, leading to corrosion. The corrosion mechanism during salt spray corrosion was electrochemical. Its development first involved anodic dissolution of the substrate:⁶⁶

Al − 3e⁻ → Al³⁺ (11)

O₂ + 2H₂O + 4e⁻ → 4OH⁻ (12)

The corrosion products adhered to the alloy surface and can provide protection. However, in chloride environments, Cl⁻ was harmful to oxide films as it can migrate to the porous layer, allowing corrosion to continue.⁶⁵ The resulting corrosion products accumulated on the alloy surface, forming a closed primary cell in certain areas, which then hydrolyzed Al³⁺ to generate H⁺.⁶⁷ The pH near the anode decreased:

Al³⁺ + 3H₂O → Al(OH)₃ + H⁺ (13)

There was a difference between the low pH solution near the anode and the high pH solution in the Cu rich area. Donatus et al.⁶⁸ indicated that the corrosion rings in Fig. 11 were caused by this pH difference.

The IMCs were difficult to dissolve in these alloys (as shown in Fig. 1). These IMCs were enriched with Cu and Mn, which were more inert than Al and therefore had a corrosion potential corrected by the Al matrix. They acted as cathodes during electrochemical corrosion, making corrosion more probable at these particles or at the interfaces between the particles and the matrix.^69–71 In the early stage of corrosion, the more active Al at the interfaces between these particles and the matrix underwent oxidation and dissolution (dealloying) from the particles or surrounding matrix. The corrosion products (porous oxides or hydroxides) generated by the dissolution of Al covered the surface of the intermetallic compound particles and matrix. At the same time, the dealloying of Al from the particles made them more enriched in Cu and Mn. Under these two effects, the particles were more inert than the Al matrix, thereby accelerating the corrosion reaction.^71,72 The gradients that existed at this time were the dealloyed area of the particles, the particles body, the surrounding matrix covered by corrosion products, and the uncorroded matrix. As the distance from the particles increased, the matrix corroded less. As the oxidation and corrosion of the surrounding matrix continued, grooves were formed near the interfaces, as shown in Fig. 19. The depth and width of the grooves gradually increased. In addition, due to the lack of physical support for the dealloying of IMCs, Cu and Mn underwent etching, dissolved into the nearby electrolyte, or deposited on the surface of the matrix. The Cu and Mn deposited on the surface of the matrix once again acted as an anode to drive the oxidation dissolution of the substrate.^73,74 This resulted in two types of corrosion pits on the alloy surface after electrochemical corrosion: (i) the dealloying of more active Al in IMCs, and the lack of support for inert Cu and Mn, which led to particles detaching (as shown in Fig. 12b–d); (ii) the corrosion pit formed by the oxidation and detachment of the Al matrix as the anode (as shown in Fig. 12e).

Fig. 19 Schematic diagram of localized corrosion in Al–Cu alloys.

Therefore, in these Al–Cu alloys, corrosion occurred first at IMCs or the interfaces between IMCs and the matrix. The uneven distribution of elements Cu and Mn due to the large number of IMCs was the reason for the slightly poorer corrosion resistance of Cu5.7.

4.3. Antibacterial mechanism analysis

Fig. 14 and 15 display the antibacterial activity of these Al–Cu alloys compared with Cu0 as the experimental control group. Antibacterial tests showed that Cu2.8 and Cu5.7 exhibited significant antibacterial activity against E. coli and S. aureus, and it was related to the Cu content.

The antibacterial mechanism of these Al–Cu alloys is discussed from a materials science standpoint in Fig. 20. The thickness of the oxide layer on the surface of an Al alloy is 4 nm so that only phases with widths larger than 4 nm penetrate the oxide film without being covered. Fig. 1 and 2 depict the microscopic morphology of the Cu-rich phases of these Al–Cu alloys. The T phases appeared as rods in the Al matrix surrounded by the θ′ phases. The width of the T phases was 20–100 nm. The width of the θ′ phases was around 10 nm, which was larger than the thickness of the oxide layer and may readily break through the oxide film, leaving the Cu-rich precipitates exposed on the surface, thereby acting as an antibacterial agent. Cu5.7, in contrast to Cu2.8, contained orthogonal needle-like θ′′ phases that were evenly distributed and were around 20 nm in length and were less than 2 nm in width. Therefore, the precipitates that can exert antibacterial effects in these alloys were mainly T phases and θ′. The volume fractions of the T phases in Cu2.8 was 3.4% and in Cu5.7 was 4.6%, indicating why Cu5.7 had better antibacterial effects. Another form of Cu was the residual IMC during the solid solution process (as shown in Fig. 1), with area fractions of 0.46% and 1.59%, respectively. In addition, a considerable portion of Cu atoms were also dissolved in solid solution in the Al matrix (Table 4).

Fig. 20 Antibacterial mechanism of Cu-containing Al alloys.

The antibacterial mechanism of these Al–Cu alloys was first considered using the antibacterial mechanism of Cu from other researchers. Cu ions, in solution some distance from the alloy surface, firmly adhere to the cell wall/membrane of bacteria, limiting their activity and range, changing their survival microenvironment, leading to metabolic turbulence, an inability to grow, and ultimately death.⁷⁵ In addition, Cu²⁺ can also penetrate the cell membranes and enter the cell, acting on the thiol groups of proteins in bacteria, causing coagulation and denaturation, resulting in enzyme inactivation, affecting DNA synthesis, and reducing bacterial reproductive ability.^76–78 Therefore, Cu²⁺ are an important reason for antibacterial effects. However, Du et al.⁷⁹ found that the concentration of Cu²⁺ in the solution was much lower than the minimum inhibitory concentration for Ti–25Cu inoculated in a suspension of S. aureus, even after 120 h. So, how does an antibacterial alloy achieve a high antibacterial rate? One suggestion is that high antibacterial rates are caused by the direct contact between bacteria and copper rich materials. Liu et al.⁸⁰ investigated the direct interaction between Ti–Cu and bacterial cell membranes using transmission electron microscopy. The direct contact between bacteria and the alloy, especially the copper rich part on the alloy surface, led to permeation of the cell membranes, which in turn facilitated the entry of Cu²⁺ into the bacterial cells. In addition, Shan Fu et al.⁸¹ proposed an antibacterial mechanism that does not rely on the toxicity of metal ions. The antibacterial effect was directly proportional to the micro potential difference (MAPD) on the alloy surface.

The above literature^75–81 indicated the antibacterial mechanism as depicted in Fig. 20 based on the microstructure of Cu2.8 and Cu5.7, the distribution of Cu, the Cu²⁺ release curves, and the outcomes of the antibacterial studies. The EDS results indicated that the mass fraction of Cu in the Cu rich IMCs was high, much higher than the Cu content in the Al matrix. This suggested that the concentration of Cu²⁺ near the Cu-rich IMCs was higher than that on the surface of the Al matrix and other parts of the solution. When the bacterial suspension came into contact with the alloy surface, corrosion caused the Cu-rich IMCs to release Cu²⁺. The Cu²⁺, dissolved from the alloys, diffused and distributed throughout the entire plane. Some bacteria came into contact with the Cu²⁺ in the solution, while others landed on the alloy surface and came into contact with the Cu-rich IMCs. The Cu-rich IMCs damaged bacterial cell membranes, and Cu²⁺ (especially high concentrations of Cu²⁺ near Cu-rich IMC) disrupted bacterial cell membranes, the expression of proteins required for survival, and affect DNA synthesis. These led to the antibacterial ability of these Al–Cu alloys.

Furthermore, the electrode potential of Al is −1.66 V, of Cu is 0.34 V, and of Mn is −1.18 V. As a result, the IMCs and the Al matrix have a potential difference. When the bacterial suspension was inoculated on the alloy surface, a micro galvanic reaction occurred between the Al matrix and the Cu-rich IMCs. The Al matrix underwent an oxidation reaction, and a reduction reaction occurred near the Cu-rich IMCs, which led to electron transfer between the Al matrix and the Cu-rich IMCs. Part of the electrons released by MAPD were transferred to the bacteria. Previous research found that some electrons transferred from the micro electric couple entered the bacteria when the bacteria came into touch with a conductive substrate.⁸² Electrons delivered into living bacteria interfered with the electron transport chain within the bacteria, resulting in bacterial death.⁸³ The potential difference between the IMCs and the Al substrate caused corrosion and dissolution, resulting in the poorer corrosion resistance of Cu5.7 compared to Cu2.8, nevertheless, the high charge transfer also interfered with normal bacterial activity and caused high antibacterial capacity.

5. Conclusions

This work designed and produced self-antibacterial Al alloys for the first time ever, giving Al alloy a new function and inspiring fresh insights into structure–function-integrated antibacterial Al alloys. The mains conclusions were as follows.

(1) Increasing the copper concentration of the antibacterial Al–xCu increased the content of copper-rich IMCs. The T phases precipitated during solid solution, while the θ series phases preferentially nucleated near the T phases during aging. Cu of the Cu2.8 was consumed during solid solution. There were no fine θ′′ phases precipitated in the matrix after aging. The Cu5.7 precipitated a large number of fine θ′′ phases in the matrix.

(2) The tiny second phases could pin dislocations and reinforce the alloy. Cu5.7 had higher UTS and YS than Cu2.8; at peak age, the UTS and YS were 348 MPa and 216 MPa for Cu5.7, respectively, which were 66% and 102% higher than for Cu2.8.

(3) Cu-rich IMCs deteriorated into matrix weak points. Increasing the IMC content decreased the corrosion resistance. The corrosion resistance of Cu5.7 was somewhat poorer than that of Cu2.8, but matched with the needed mechanical qualities and corrosion resistance for typical Al–Cu alloy applications.

(4) Al–xCu alloys showed good antibacterial properties. After 24 h, Cu5.7 virtually eradicated all S. aureus. The potential of Al alloys incorporating Cu as antibacterial structure–function integrated materials was confirmed by this investigation.

Data availability

All data included in this study are available upon request by contacting the corresponding author.

Conflicts of interest

All authors agree to this submission and declare no interest conflict.

Acknowledgements

Financially supported by the Shandong Province Natural Science Foundation (No. ZR2023ME181), the National Natural Science Foundation of China (No. 52305313) and the Natural Science Foundation of Hunan Province (No. 2023JJ40553 and No. 2023JJ60433).

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Footnote

† These two authors contributed equally.

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