1. Introduction
The 7000 series alloys based on the Al-Zn-Mg-(Cu) system are used for high-strength structural components in aerospace, automotive, and military applications. These precipitation-hardenable alloys exhibit tensile yield strengths approaching 600 MPa due to densely distributed nano-sized precipitates formed during artificial aging [
1].
The precipitation sequence in 7000 series alloys—like other precipitation-strengthened alloys—is influenced by alloy chemistry, thermo-mechanical processing, and final age hardening heat treatments. The Zn:Mg ratio, Cu content, homogenization and rolling practice, and aging practice collectively determine the final volume fraction and spatial distribution of precipitates in Al-Zn-Mg-(Cu) alloys [
2].
Precipitates typically observed in artificially aged 7000 series alloys without Cu (Al-Zn-Mg) include the equilibrium
η (MgZn
2) phase and its precursor
η’ phase. In artificially aged 7000 series alloys with Cu (Al-Zn-Mg-Cu), observed precipitates are the equilibrium
η phase expressed as Mg(Zn,Cu,Al)
2 and its precursor
η’ phase [
3]. In general, the precipitation sequence in Al-Zn-Mg-(Cu) alloys begins with the decomposition of a supersaturated solid solution (SSSS) into nano-sized (~3 nm) clusters of Mg and Zn atoms called Guinier-Preston (GP) zones. Two types of GP zones can form depending on quenching and aging conditions: spherical GPI zones or plate-like GPII zones [
4]; both types impede dislocation movement and thus increase strength. GP zones evolve into the metastable ellipsoidal
η’ strengthening phase that is semi-coherent with the aluminum matrix [
2,
5,
6]. Upon further aging, the
η’ evolves into the equilibrium
η phase [
6].
Precipitation in metallic alloys causes local composition fluctuations as precipitates nucleate and grow. Small-angle X-ray scattering (SAXS) signals are sensitive to local changes in electron density or atomic number, and thus SAXS is sensitive to local changes in composition and can be used to study precipitate evolution.
High-energy synchrotron X-ray sources enable SAXS experiments on metallic alloys. For sheet materials, SAXS experiments can be run in transmission mode, where the incident X-ray beam passes through the sample. The scattered X-ray signal can be analyzed to determine the precipitate size and volume fraction [
7,
8]. The scattered signal or intensity
I is measured as a function of the scattering vector
q.
For synchrotron experiments, the X-ray wavelength λ is often fixed. In SAXS experiments, the scattering angle 2θ can range from 0.1 to 6°. The scattering angle is determined by the distance between the sample and the detector as well as the detector size and the beam stop position. In transmission mode, the scattered intensity is recorded as a two-dimensional image. The two-dimensional scattering image is radially averaged, producing a one-dimensional scattering curve of intensity I vs. scattering vector q.
The measured scattering signal
I(q) is proportional to the squared difference between the scattering length densities of the scatterer
ρe,scatterer, and matrix
ρe,matrix.
Scattering length density is related to atomic number. SAXS is most effective when the difference between the atomic number Z of a scatterer and the atomic number of the matrix is large. The study of precipitation in 7000 series alloys is well suited for SAXS because there is a high electron density contrast between zinc-bearing precipitates (ZZn = 30) and the aluminum (ZAl = 13) matrix.
The average precipitate size and volume fraction can be extracted from the scattering curves. The precipitate size and volume fraction information can be used to predict the strength contributed by precipitation hardening.
3. Results and Discussion
Hardness and electrical conductivity vs. heat treatment conditions are displayed in
Figure 4 for the Al-Zn-Mg and Al-Zn-Mg-Cu alloys. Both alloys reached a peak hardness after 3 h at 140 °C (
Figure 4a). The copper-bearing alloy had a greater peak hardness after 140 °C/3 h compared to the non-Cu-bearing alloy (90 vs. 85 HRb). For the Cu-bearing alloy only, the hardness appeared to plateau with very little change between the 120, 140, and 160 °C isothermal heat treatments.
As-quenched hardness (13 HRb) for the Al-Zn-Mg alloy (
Figure 4a) was lower than the as-quenched hardness (57 HRb) for the Al-Zn-Mg-Cu alloy (
Figure 4a). After 24 h of natural aging, the Al-Zn-Mg alloy gained considerable strength, as indicated by the sharp increase in hardness from 13 to 52 HRb (
Figure 4a). Chinh et al. [
18] concluded that Cu-bearing vacancy-rich clusters (VRCs) can form immediately after quenching, offering a significantly greater strengthening effect than VRCs formed in ternary Al-Zn-Mg alloys. These Cu-bearing VRCs may explain the large difference in the observed as-quenched hardness since more VRCs would result in more GP zones and higher strength.
In both alloys, regardless of copper content, natural aging after quench resulted in increased hardness and decreased electrical conductivity (
Figure 4b). This confirmed GP zone formation, as GP zones impede dislocation movement and are thought to impair lattice periodicity, resulting in more restrictive electron movement, and thus reduced conductivity [
19]. In contrast, conductivity increased with increasing heat treatment temperature for both Cu- and non-Cu-containing alloys. The increase in conductivity was due to the decomposition of the solid solution into precipitates. Solute in solid solutions tends to restrict electron movement. As solutes leave solid solution and form precipitates, electrons tend to move more freely throughout the aluminum matrix, resulting in increased electrical conductivity.
Figure 5 shows TEM images of the two alloys after the 120 °C and 160 °C 3 h isothermal heat treatments. After the 160 °C/3 h treatment,
η’/η phases were observed in the aluminum matrix for both the Al-Zn-Mg (
Figure 5a) and Al-Zn-Mg-Cu (
Figure 5b) alloys. Precipitates in the Al-Zn-Mg alloy after the 120 °C/3 h treatment (
Figure 5c) had little contrast, making observation difficult. However, the HRTEM inset image in
Figure 5c shows evidence of coherent GP zones as dark agglomerates, with similar lattice structure to the surrounding light-gray aluminum matrix. Precipitates can be clearly observed in the TEM image of the Al-Zn-Mg-Cu alloy after the 120 °C/3 h treatment (
Figure 5d). The HRTEM inset shows that these precipitates were coherent GP zones, indicated by the similarity in lattice structure between the dark contrast areas and light-gray aluminum matrix (
Figure 5d). In summary, TEM observations indicated that GP zones were present after the 120 °C/3 h heat treatment for both the non-Cu- and Cu-containing alloys. After the 160 °C/3 h treatment,
η’/η precipitates were observed in both alloys.
The scattering curves for each 3 h isothermal heat treatment are plotted for the Al-Zn-Mg and Al-Zn-Mg-Cu alloys in
Figure 6. The scattering curves from the Al-Zn-Mg-Cu alloy are plotted in
Figure 6a,b, for 100, 120, and 140 °C isothermal heat treatments (
Figure 6a) and 160, 180, and 200 °C heat treatments (
Figure 6b). The scattering curves for the non-Cu-containing Al-Zn-Mg alloy are plotted in
Figure 6c,d for 100, 120, and 140 °C isothermal heat treatments (
Figure 6c) and 160, 180, and 200 °C heat treatments (
Figure 6d). The scattered intensity at the high
q-range is due to small precipitates such as GP zones and small, early-stage
η’ precipitates. Scattered intensity at the low
q-range is due to larger precipitates such as
η’ and
η phases. As isothermal heat treatment temperature increased, the curves shifted to lower
q-range values and higher intensities as the precipitate size and volume fraction increased.
For the scattering curves at lower temperatures (e.g., 100 °C, 120 °C, 140 °C), the curves begin with a sharp decline in intensity at low
q, then rise to a maximum, followed by a gradual decline in intensity. This initial dip in intensity at low
q is caused by a destructive interference effect due to high precipitate number densities [
8]. This initial intensity dip was filtered out prior to modeling the precipitate size distribution using the GSAS-II maximum entropy method. The red dashed lines overlaid on the scattering curves in
Figure 6 represent the portion of the scattering curve that was modeled using the maximum entropy method. The particle size distributions were calculated from these best-fit functions.
The average precipitate diameter determined from the SAXS-MEM is plotted for each artificial age condition for the non-Cu- and Cu-containing alloys in
Figure 7. The solid points plotted in
Figure 7 are the average precipitate diameters measured from the TEM images, which are in good agreement with the SAXS measurements. The calculated precipitate volume fraction is shown in
Figure 8. The volume fraction for the Cu-containing alloy was higher than the non-Cu alloy at low heat treatment temperatures (i.e., 100 °C, 120 °C). After the 140 °C and 160 °C heat treatments, both alloys had nearly the same volume fraction. The volume fractions in both alloys plateaued after the 180 °C and 200 °C heat treatments. The Cu-containing alloy had about a 20% higher volume fraction than the non-Cu alloy after the 180 °C and 200 °C heat treatments.
Figure 9 shows the calculated strength increase due to precipitation hardening for both pure shear and pure by-pass mechanisms. The strength increase was calculated for each case using Equations (6) and (7); the precipitate size and volume fraction measurements presented in
Figure 6 and
Figure 7 were used as inputs. The strength of a precipitation-hardenable alloy is governed by the interaction of the dislocations with the precipitates. Dislocations interact with precipitates by two mechanisms: (1) shearing or (2) by-pass. In the under-aged condition, the shearing mechanism is dominant—where strength increase due to precipitation hardening Δσ is proportional to the precipitate volume fraction
fV and average precipitate radius
R,
Here, precipitate size is proportional to strength. Precipitates are shearable in the under-aged condition up to a critical radius. When precipitate size grows beyond the critical radius, then the Orowan strengthening mechanism becomes dominant. Instead of shearing, the dislocations bow around and by-pass precipitates. This is called Orowan strengthening, where
is proportional to
fV and
R by:
Precipitates in the over-aged condition are non-shearable, and strength is controlled by the Orowan mechanism, where material strength is inversely proportional to precipitate size.
The transition from shear to by-pass mechanism occurs around the 140 °C heat treatment temperature for both Al-Zn-Mg and Al-Zn-Mg-Cu alloys. Precipitate shearing is the dominant strengthening mechanism after the 100 and 120 °C heat treatments, whereas the by-pass mechanism becomes dominant after the 160, 180, and 200 °C heat treatments.
Peak strength in precipitation-hardenable alloys occurs at the transition from shearing to by-pass. Peak hardness (strength) was observed after the 140 °C/3 h heat treatment for the Cu-bearing and non-Cu-bearing alloys in
Figure 4a. The point at which peak hardness is observed in
Figure 4a agrees well with the calculations in
Figure 9, which shows that the transition from shearing to by-pass mechanism occurs around 140 °C.
The average precipitate diameter measured from the SAXS data for the 140 °C/3 h heat treatment was 44 ± 4 Å for the non-Cu alloy and 52 ± 5 Å for the Cu-containing alloy. Assuming these measurements are respective of critical precipitate size, the transition from shearing to Orowan type strengthening mechanism occurred at larger precipitate sizes in the Cu-containing Al-Zn-Mg-Cu alloy. Hardness began to decrease after the 160 °C aging treatment for both alloys, followed by further decline after the 180 °C and 200 °C heat treatments (
Figure 4a). Similarly, the calculated strength delta in
Figure 9 decreases after the 160 °C heat treatment. As the hardness decreases, precipitates continue to coarsen, indicated by an increase in average precipitate size for both alloys in
Figure 7.