The effects of natural ageing treatment prior to artificial ageing treatment on the microstructures and mechanical properties of AA7075 Al-5.7Zn-2.6Mg-1.5Cu-0.18Cr-0.08Mn-0.05Si-0.17Fe (wt.%) aluminum alloy have been investigated. The hardness of solution-treated samples (91.0 HV) profoundly increased to 146.8 HV after 7 days of natural ageing. The purpose of the present work was to examine the kinetic hardening evolution in subsequent artificial ageing treatments of samples naturally aged for 7 days and their counterparts without natural ageing. The former were labelled as NA-7d samples, and the latter, NA-0d samples. After artificial ageing at 120 °C for 2 h, the hardness of NA-0d samples increased rapidly to 148.2 HV, which was approximately the same as that of the specimens with natural ageing for 7 days, compensating for the prior state of lower hardness without natural ageing. After being treated at 120 °C for 16 h, the ultimate tensile strength (UTS) and yield strength (YS) of NA-7d reached the highest value, respectively, 601 MPa and 539 MPa, followed by a slight decrement of UTS when aged to 24 h. On the other hand, NA-0d specimens aged at 120 °C for 16 and 24 h showed nearly the same UTS (598 MPa); the former possessed YS of 538 MPa, and the latter, 545 MPa. The results presumably reveal that the peak ageing condition for NA-0d samples can be achieved under 24 h ageing at 120 °C. Under the same treatment at 120 °C for 24 h, the size of η’ phase in NA-7d sample (with a length of 4.96 nm) coarsened and grew larger than that in NA-0d sample (with a length of 3.46 nm). In addition, some η’ phase in the NA-7d sample was found to be transformed into the η2 phase. The results indicated that the naturally aged specimens (NA-7d) reached the peak ageing condition earlier, but did not significantly enhance the UTS in AA7075 aluminum alloy, as compared to the samples without prior natural ageing (NA-0d).
A nitrogen DC-pulse atmosphere pressure plasma jet (APPJ) is used to convert ferric nitrate (Fe(NO3)3) and chloroplatinic acid (H2PtCl6) mixed liquid precursor films into PtFe nanocompounds on a fluorine-doped tin oxide (FTO) substrate. Scanning transmission electron microscopy indicates nanoparticles distributed on a thin continuous layer on the FTO substrate. The APPJ-synthesized PtFe nanocompounds contain a mixture of crystalline and amorphous phases. X-ray photoelectron spectroscopy shows that most Pt is in the metallic phase and most Fe, in the oxidized phase. A dye-sensitized solar cell (DSSC) with only 5-s APPJ-processed PtFe counter electrode (CE) shows significantly improved efficiency. This suggests the rapid processing capability of the nitrogen DC-pulse APPJ. A PtFe prepared with higher H2PtCl6/Fe(NO3)3 volume ratio shows better catalytic performance, as confirmed by cyclic voltammetry, electrochemical impedance spectroscopy, and Tafel experiments. The DSSC with APPJ-processed PtFe CE shows comparable efficiency to that of 15-min furnace-calcined Pt CE, suggesting that the APPJ processed PtFe requires less Pt.
Cs-corrected high-angle-annular-dark-field scanning-transmission-electron microscopy (Cs-corrected HAADF-STEM) was employed to examine the phases in AlCuLi alloy (AA2050), including GP(T1), GP(θ″) and GPB zones with their subsequent nanometer-sized products, T1 (Al2CuLi), θ′ (Al2Cu), and S (Al2CuMg) precipitates, respectively. Under the peak-aging condition, some solute-atom enriched clusters could still be found, and the newly-formed nucleus of GP(θ″) with a mono-layer {100} plane of Cu atoms occurred at the adjacent area of the joint between θ′ and S precipitates or the edge of an individual S precipitate. The transition of a single Cu-layer GP(θ″) → θ′ was presumed to be transformation via in-situ nucleation. The developing GP(θ″) zones and θ′ precipitates were easily subjected to soft impingement. However, hard impingement between two variants of θ′ presumably occurred, wherein one θ′ variant precipitate was blocked out by the other θ′ variant. As for the creep-ageing forming (CAF) treated sample, some precipitates of T1 and θ′ were found to have the cutting characteristic on specific ledges.
Sn is a well-known grain boundary segregation element that improves the machinability of steel. Sn has been considered as a replacement for Pb in super-free-machining steels. The effect of Sn on the microstructure of the Fe-0.05C-0.03Si-1.28Mn-0.36S-0.05P base composition containing 0.0002 (without addition), 0.062, 0.12, and 0.18 Sn (wt.%) low-carbon free-machining steels was investigated though thermodynamic calculations, optical microscopy, scanning electron microscopy, transmission electron microscopy, electron probe microanalysis, and high-temperature laser scanning confocal microscopy. The microstructure of the free-machining steels was composed of α-ferrite, pearlite, and manganese sulfide. Sn significantly decreased the pearlite content of the steels. Most of the Sn was dissolved in the matrix, and the remainder was dissolved in manganese sulfide. No FeSn intermetallic compound precipitation was observed through transmission electron microscopy, but a significant strengthening effect was observed in α-ferrite. Sn had little effect on the solidification behavior or sulfide precipitation behavior of the steels when its content was lower than 0.2 wt.%. Similar to Al and Si, Sn is a ferrite-stabilizing element that expands both the δ-ferrite and α-ferrite phase regions, promotes α-ferrite formation, and inhibits carbide precipitation. Sn segregates at the interface, decreasing the interfacial energy and promoting the Widmanstätten ferrite phase transformation. A much lower cooling rate than that of Sn-free free-machining steels should be adopted to restrain the formation of Widmanstätten ferrite after hot rolling. Sn addition greatly improved the machinability of the experimental steels.
The microstructures and mechanical properties in the peak ageing condition, warm forming process, and paint baking process of two Al–Zn–Mg–Cu aluminum alloys, which possessed the same base compositions but different levels of Mn (0 and 0.24 wt%), have been investigated. The results show that Mn addition can hardly refine the grain size (around 20 μm) in AA7075 aluminum alloy in warm forming and paint baking processes. Mn-free and Mn-containing alloys subjected to all three processes possess nano-precipitates of nearly the same size, indicating that Mn addition does not influence the nano-precipitate evolution in the aluminum matrix. Energy-dispersive X-ray spectroscopy (EDS) mapping results of dispersoids show that dispersoids in Mn-bearing specimens are enriched in Mn and depleted in Zn and Mg. The Zn and Mg expelled from dispersoids in Mn-containing specimens into the aluminum matrix presumably formed Zn- and Mg-rich zones around the dispersoids, which improved the precipitate transformation reation from GPII to η′ phase at the boundaries between aluminum matrix and dispersoids. Although the solid solution strengthening effect triggered by Mn addition might be less effective because of the formation of dispersoids, the slight enhancement of ultimate tensile strength (UTS) in Mn-containing specimens rather than in Mn-free specimens in all states can be attributed to the formation of nano-precipitates around the dispersoids.
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