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Article

The Microstructure and Properties of Al–Mn–Cu–Zr Alloy after High-Energy Ball Milling and Hot-Press Sintering

by
Olga A. Yakovtseva
,
Andrey G. Mochugovskiy
,
Alexey S. Prosviryakov
,
Andrey I. Bazlov
,
Nadezhda B. Emelina
and
Anastasia V. Mikhaylovskaya
*
Department of Physical Metallurgy of Non-Ferrous Metals, National University of Science and Technology “MISIS”, Leninskiy pr. 4, Moscow 119049, Russia
*
Author to whom correspondence should be addressed.
Metals 2024, 14(3), 310; https://doi.org/10.3390/met14030310
Submission received: 6 February 2024 / Revised: 28 February 2024 / Accepted: 3 March 2024 / Published: 6 March 2024
Browse Figures
Versions Notes

Abstract

:
In the present research an Al–7.7%Mn–4.9%Zr–3.2%Cu (wt%) alloy was processed by mechanical alloying (MA) followed by hot press sintering. The microstructure, phase composition, and mechanical properties of the MA granules and sintered samples were investigated. The dissolution of Mn, Zr, and Cu with further precipitation of the Al6Mn phase were observed during high-energy ball milling. In the alloy processed without stearic acid after milling for ~10 h, an Al-based solid solution with ~4.9 wt%Zr, ~3.2 wt%Cu and a ~5 wt%Mn with a grain size of ~16 nm and a microhardness of ~530 HV were observed. The addition of stearic acid facilitated Mn dissolution and precipitation of the Al6Mn phase during milling but led to the formation of the ZrH2 phase that decreased the Zr solute and the microhardness. Precipitation of the Al6Mn, L12–Al3Zr, and Al2Cu phases during annealing and sintering of the MA granules in the temperate range of 350–375 °C was observed, and an additional Al20Cu2Mn3 phase was precipitated at 400–450 °C. Hot-press sintering at 450 °C provided a low fraction of cavities of ~1.5%, the yield strength of 1100 MPa, ultimate compressive strength of 1200 MPa, strain at fracture of 0.5% at room temperature, the yield strength of 380 MPa, ultimate compressive strength of 440 MPa, and strain at fracture of 3.5% at 350 °C. The microstructural evolution during high-temperature deformation on the sample surface was studied and the differences in deformation behavior for the alloys sintered at different temperatures were discussed.

Graphical Abstract

1. Introduction

Aluminum-based alloys are frequently used in the aircraft and automotive industries due to their attractive combination of properties, such as low density, good corrosion resistance, and mechanical properties [1,2]. Aluminum is the second-most widely used metal after iron. About 100 million tons of Al and its alloys are produced per year for different applications, which is ~50% of all industrial metals with the exception of iron [3]. The Al market is growing every year due to the increasing use of Al in the transportation industry. Conventional wrought and cast aluminum-based alloys are not able to meet the industry’s requirements for a combination of properties [4,5]. High-strength light metallic alloys and composite materials are required for operating at high loads and elevated temperatures [6,7,8,9,10,11]. Due to a good combination of performance, economic, and processing properties, Al–matrix composites account for around 60–70% of all metal–matrix composites [12,13,14,15,16,17,18]. The development of aluminum-based alloys and Al–matrix composites with increased strength is an actual scientific issue that is dictated by relevant industrial requirements.
A traditional approach by which to increase the strength properties of aluminum alloys is based on the precipitation strengthening effect. Alloying with Mn and Zr provides precipitation of fine dispersoids that are stable at elevated temperatures [19,20,21,22,23]. Manganese and zirconium have a low equilibrium solubility in aluminum, while a supersaturated Al-based solid solution forms due to rapid solidification [24,25,26,27]. The subsequent decomposition of the solid solution leads to the formation of nanoscale intermetallic particles (dispersoids), which significantly improve the yield strength [28,29,30,31,32,33,34]. During decomposition of the solid solution in a temperature range of 350–450 °C, an Al3Zr phase with L12 crystal structure and nanoscale size of ~10 nm is formed [35,36]. At the higher temperatures, coarser particles of the D022 or D023 Al3Zr phases and a weaker strengthening effect are observed [37,38,39]. The equilibrium phase for Al–Mn-based alloys is orthorhombic Al6Mn, which is formed during both solidification and solid solution decomposition at elevated temperatures. Ageing in a temperature range of 300–400 °C leads to the precipitation of the icosahedral I-phase and the Al12Mn phases [40,41,42,43]. The precipitation strengthening effect is noticeable for alloys with both Zr and Mn [44,45,46,47,48,49,50]. The strengthening capacity and ageing effect of the elements can be improved by increasing their solubility in Al.
Mechanical alloying (MA) can increase solid solubility [51,52,53,54]. MA is a powder technology that is based on welding, fracturing, and rewelding powder particles during the high-energy ball milling process [55,56,57,58,59]. The severe plastic deformation accumulated during MA provides the nano-grained structure with high dislocation density and vacancy concentration [60]. Large concentrations of the lattice defects are conductive to high diffusivity, which increases the solubility of the alloying elements. The abnormally supersaturated solid solutions are formed during MA of Al–Mg [61,62], Al–V [63,64], Al–Fe [65,66], Al–Mn [67,68,69,70], and Al–Zr [71,72,73] alloys. The solubility of Zr and Mn is increased to ~8 at% [71] and ~3.1–3.5 at% [67,74], respectively. Dissolution kinetics is a diffusion controlled process that depends on grain refinement and the accumulation of lattice defects [75,76,77,78,79,80]. Therefore, alloy chemical composition, milling parameters, rotation speed, and the addition process control agent (PCA) [74,81,82,83] can all influence the solute content and the final properties of the alloys.
To process bulk nanocrystalline materials, the mechanically alloyed powders are subjected to sintering at elevated temperatures [84,85,86]. The most effective sintering methods are the hot-press sintering [87,88], hot extrusion [89,90], spark plasma sintering (SPS) [66,91], and microwave sintering [92]. These methods are actively used to produce ceramic particle-reinforced aluminum–matrix composites [51,93]. Sintering at elevated temperatures results in decomposition of the supersaturated solid solution in the MA granules. The influence of sintering temperature on the decomposition of the supersaturated solid solution of MA alloys is poorly understood. Several studies have demonstrated the high strength and heat resistance of bulk alloys produced by mechanical alloying of aluminum with several transition metals. High-energy ball milled and hot-press sintered Al-based alloys with 5 at% Fe [66] and 3 at% Zr [73] exhibit high compressive strengths of 800–1000 MPa at room temperature and 350–500 MPa at 300–350 °C. The increase in strength was achieved due to the nanograined structure and nanoscale precipitates, which provide Hall–Petch [94,95] and Orowan strengthening mechanisms, respectively.
In the present work, the microstructure and properties of Al–Mn–Zr–Cu alloys processed by mechanical alloying followed by hot-press sintering have been investigated. The influence of the milling time and the effect of a PCA (stearic acid) addition on the evolution of the phase composition, microstructural parameters and microhardness during mechanical alloying and the effect of the sintering temperature on the compressive mechanical properties for the bulk samples were investigated and discussed.

2. Materials and Methods

The Al–Mn–Zr–Cu alloy was prepared by mechanical alloying of the mixed chips of the Al–14.3 %Mn–6.5 %Cu alloy (as-cast microstructure of the master alloy having the aluminum solid solution, while the Al6Mn and Al20Cu2Mn3 phases have been described in [74]) and the 99.9 Zr and 99.7 Al powders (wt%). Alloy was processed using a Retsch PM400 planetary ball mill (VERDER Company, Haan, Germany) with a rotation speed of 300 rpm. Stainless steel balls (AISI 52100 grade) were used. The ball-to-alloy weight ratio was 20:1. The milling was processed for 5 min and stopped for 5 min to prevent overheating. The applied milling parameters provided a high efficiency of the milling process for Al-based alloys [69,71,72,73,74]. Mechanical alloying was carried out in an argon atmosphere to prevent oxidation. To study the effect of the process control agent and to decrease a potential cold welding effect, the milling was processed without and with ~1 wt% stearic acid [96]. The nominal composition of the alloy was Al–7.7 %Mn–5 %Zr–3.5 %Cu. The alloying elements content was controlled for PCA-free alloy via spectral analysis in an inductively coupled plasma-atomic emission spectrometer CAP-6300 (Thermo, Waltham, MA, USA) after milling for 10 h. The actual composition of the alloy was Al–7.7%Mn–4.9%Zr–3.2%Cu. The content of residual Fe was ~0.35%, and it was introduced from the milling media.
The mechanically alloyed PCA-free granules processed for 10 h were annealed to determine the potential influence of the hot-press sintering on the alloy phase composition. To prevent oxidation, the annealing was carried out in vacuum quartz tube at temperatures in a range of 350 to 450 °C for a time range of 0.5 to 8 h.
Hot-press sintering of the 10-h-milled granules was carried out at a load of 0.5 GPa using a Zwick Z250 universal testing machine at the temperatures of 350, 375, 400, 450 °C. Bulk samples with a diameter of 6 mm and a height of 9 mm were obtained and used for further studies of density and mechanical properties. The granules were pre-pressed in a steel matrix at room temperature then heated to a required temperature and compressed in an isothermal condition, kept under the load for 5–7 min in order to stabilize the stress. To decrease the friction effect, a graphite-based lubricant was used.
The mean microhardness values of the granules before and after annealing and sintering were determined using a 402MVD microhardness tester (Wilson & Wolpert, Fort Worth, TX, USA) by a Vickers method at a load of 25 g using 15–20 measurements. Compression tests of the sintered samples with a diameter of 6 mm and height of 9 mm were carried out using a Zwick Z250 universal testing machine at room temperature and at the temperature of 350 °C. The initial strain rate was 10−3 s−1. Three samples per point were tested and a mean value was calculated.
X-ray diffraction patterns (XRD) were achieved using a D8 Discover diffractometer Bruker-AXS (Bruker Corporation, Billerica, MA, USA) with Cu–Kα radiation after the mechanical alloying for different times, annealing of MA granules, and hot-press sintering of the samples. XRD patterns were taken in a wide-angle range of 2θ of 20 to 140° with a step size of 0.05°. The values of the crystallite sizes (coherent scattering region (CSR)) were calculated using the Williamson–Hall method [97]. The lattice parameter was estimated via extrapolation method [71]. The error bars for the Al lattice parameter were <0.0001 nm.
Differential scanning calorimetry (DSC) was performed using a Setaram Labsys DTA/DSC 1600 calorimeter and heating within a temperature range of 20–680 °C with a heating rate of 20 K/min. The weight of the granules for DSC study was ~80 mg.
Microstructural analysis was performed using a JEOL-JEM 2100 (JEOL, Tokyo, Japan) transmission electron microscope (TEM) and a TESCAN VEGA 3LMH scanning electron microscope (SEM) (Tescan Brno s.r.o., Kohoutovice, Czech Republic) equipped with an X-MAX80 energy-dispersive X-ray spectrometer (EDS) (Oxford Instruments, Abingdon, UK). The size of the granules was estimated as the mean diameter. The grain size was calculated using dark field TEM images via liner secant method. The samples preparation for the microstructural studies after sintering included grinding on SiC papers of various dispersions, polishing with chromium oxide and final polishing with OPU-suspension (Struers APS, Ballerup, Denmark).

3. Results

3.1. Mechanical Alloying

The mixture of components was subjected to mechanical alloying for up to 15 h. Diffraction peaks for Zr and Cu-containing phases were not observed after milling for 5–15 h for samples both with and without PCA (Figure 1). According to the diffraction patterns, the ZrH2 phase was formed after milling for 5–15 h, indicating that the stearic acid led to the formation of zirconium hydride (Figure 1a). There were no Zr and Zr-bearing phases for the PCA-free sample after milling for 7.5–12.5 h. When increasing the milling time to 10 h for the alloy with PCA and to 12.5 h for the PCA-free alloy, the relative intensity of the Al6Mn diffraction peaks increased. The maximum relative intensity of the Al6Mn phase peaks was observed after milling for 15 h for the alloy with stearic acid.
For the PCA-free alloy, fine granules with a size of ~1–10 µm and their agglomerates with a size of ~100–300 µm were observed after milling for 5 h (Figure 2a). The mean granular size was 52 ± 6 µm. The alloy with PCA showed fine granules with a size of about 1 µm and a small fraction of agglomerates with a size of ~10–50 µm; the mean granular size was 3 ± 1 µm (Figure 2c). With an increase in milling time from 5 to 10 h, the granular size for the PCA-free alloy decreased to 12 ± 1 µm (Figure 2b). The granular size did not change significantly for the alloy with PCA; a mean size was 5 ± 1 µm after milling for 10 h (Figure 2d). SEM-EDS analysis of the alloys after milling for 5 h confirmed the uniform distribution of the alloying elements Mn, Zr, and Cu for the PCA-free alloy (Figure 2e). Meanwhile, Zr-enriched areas were observed for the alloy with PCA (Zr distribution map in Figure 2f). According to the XRD data, these areas should belong to ZrH2 (H could not be defined by EDS). Residual Fe-rich inclusions were introduced from the milling balls (see Fe distribution maps in Figure 2e,f).
During milling, Al diffraction peaks shifted from their equilibrium position (Figure 1c), indicating the change in lattice parameter. The Al lattice parameter changed with a minimum for alloys processed both with and without PCA (Figure 3a). The decrease in the lattice parameter can be associated with Mn and Cu dissolution, but an increase in its value can be associated with Zr dissolution and/or precipitation of the Mn- and Cu-bearing phase. The XRD data suggest that Mn dissolution occurred during milling for 5 h for the alloy with PCA and during milling for 10 h for the alloy without PCA. Increasing the milling time resulted in the precipitation of the secondary Al6Mn phase. Meanwhile, the Al6Mn phase of solidification origin was incompletely dissolved. Zirconium was dissolved in Al during the early stages of milling for 5 h for the PCA-free alloy, and it formed ZrH2 for the alloy with stearic acid.
Considering the error bars, the mean crystallite size (coherent scattering region) changed insignificantly during mechanical alloying. The mean CSR value was in a range of 20–24 nm for the alloy with PCA and 16–22 nm for the alloy without PCA.
Microhardness increased with increasing milling time (Figure 3c). For the alloy with PCA, the hardness changed from ~250–300 HV after milling for 2.5–10 h to ~450 HV after milling for 15 h. For the PCA-free alloy, hardness was significantly higher, increasing from 450 HV after milling for 5 h to 625 HV after milling for 12.5 h.
TEM studies of the alloy milled for 10 h confirmed its nanocrystalline structure (Figure 4). SAED showed a ring-type pattern corresponding to Al grains surrounded by high-angle grain boundaries. The mean grain size of Al solid solution measured from dark-field images was 16 ± 4 nm, which was similar to the CSR size determined by XRD. Some scattering reflexes for particles of the second phase with larger interlayer distances were also observed in the SAED (see arrows in Figure 4c). According to the XRD data, these reflections may belong to the Al6Mn phase with planes (200) and (113). The PCA-free alloy milled for 10 h with high hardness and a high Zr solute content was used for further annealing and hot-press sintering.

3.2. Annealing of the MA Granules

The recovery, recrystallization, grain growth, and decomposition of the supersaturated solid solution with the precipitation of secondary phases are possible during heating and hot-press sintering of the MA granules. In order to analyze these processes and select an appropriate sintering time, thermal analysis, hardness and XRD data evolution were investigated for MA granules after annealing in a temperature range of 350–450 °C.
Four exothermic peaks (P1, P2, P3, P4) and one endothermic peak (P5) were observed in the DSC curve (Figure 5). The first exothermic peak P1 was in a temperature range of 140–250 °C. This peak was non-symmetric and suggested that two effects occurred in this temperature range. The second P2 peak was higher and appeared in a temperature range of 275–350 °C with the maximum at ~300 °C. Two small exothermic peaks were also observed at higher temperatures of ~450 °C and ~530 °C. The exothermic peaks exhibited the following reaction enthalpies: ~1 J/g for P1, ~4 J/g for P2, and ~0.5 J/g for both P3 and P4 peaks. The P5 endothermic peak corresponds to the melting of the studied alloy and indicates a reaction enthalpy of about −30 J/g.
The PCA-free granules milled for 10 h were further annealed in a temperature range of 350–450 °C for a time range of 0.5–32 h and were water cooled (Figure 6) to analyze the evolution of hardness and phase composition of the granules at elevated temperatures. A similar behavior with a hardness minimum of ~450–480 HV after annealing for 0.5–1 h and 2 h was observed at the temperatures of 350 and 400 °C, respectively. The hardness decreased insignificantly, to 490 HV, after annealing for 0.5 h at a temperature of 450 °C, but longer annealing for a time of 1–8 h resulted in a significant decrease in hardness to 380–400 HV.
The XRD data of the granules annealed at temperatures of 350–450 °C for 2–4 h are presented in Figure 7. The phase composition of the granules changed during annealing. The precipitation of the Al6Mn phase was clearly observed after annealing in the studied temperature range. The Al3Zr, Al2Cu phases were defined after annealing but key peaks of the highest intensity of these phases overlapped with peaks of the other phases. The Al20Cu2Mn3 phase was revealed after annealing at 400 and 450 °C. The positions of Al peaks were shifted to the left after annealing, and the lattice parameter increased from 0.4039 to 0.4048–0.4055 nm (Table 1). Thus, the decomposition of the Al solid solution occurred, and the solute content of Mn and Cu decreased. Meanwhile, the lattice parameter was higher than that of for Al, indicating incomplete decomposition of Al solid solution. The mean crystallite size increased to ~25–30 nm after annealing at 350–450 °C.

3.3. Hot-Press Sintering

The granules were subjected to hot-press sintering at 350, 375, 400, and 450 °C. The microstructures of the samples after sintering are shown in Figure 8. The hardness, density and compression test data for the hot pressed and further annealed sintered samples are summarized in Table 2.
The samples pressed at 350 °C demonstrated a high fraction of cavities of about 12 ± 4% and a low density of 2.70 ± 0.03 g/sm3. The density of the samples changed insignificantly after annealing for 2 h at a temperature of 350 °C.
The density of the samples pressed at 375–400 °C was ~2.77–2.78 g/cm3. Further annealing for 2 h at the sintering temperatures resulted in a higher density of 2.85–2.89 g/cm3 and a cavities fraction of ~6 ± 1%. The maximum density of ~3.05 ± 0.02 g/cm3 was obtained after hot pressing at 450 °C. Further annealing at 400 °C for 2 h increased the density of the samples to 3.15 ± 0.05 g/cm3. This treatment provided a low fraction of cavities of ~1.5 ± 0.5%. The annealing temperature of 400 °C was lower than the hot-pressing temperature of 450 °C, this was in order to avoid strong softening during the long-term annealing at 450 °C (see HV–time curve in Figure 6).
The microhardness was 510–520 HV for the samples sintered at 350 and 375 °C (Table 2), which was similar to the MA granules of 530 ± 10 HV. During hot pressing, the heating and holding time was ~15 min. The short time treatment at elevated temperature insignificantly decreased the hardness of the MA granules (Figure 6) and provided a comparatively high hardness for the sintered samples. Lower values of 470 ± 20 HV and 380 ± 20 HV were obtained after sintering at 400 and 450 °C, respectively. The hardness of the sintered samples changed insignificantly (within the error bars) after annealing for 2 h.
XRD analysis of the sintered samples revealed two main phases of Al and Al6Mn (Figure 9). The Al20Cu2Mn3 phase was also observed after sintering at 400 and 450 °C. The Al3Zr phase was defined after sintering but key peaks of the highest intensity overlapped with Al peaks. The broadening of Al peaks decreased and an intensity for the Al6Mn phase peaks increased with increasing temperature. The mean crystallite size for the sintered samples varied in a range 25–30 nm, and these values did not depend on the sintering temperature. It is notable that similar values were obtained for the annealed granules. The Al lattice parameters were 0.4042, 0.4043, 0.4047, 0.4053 nm after sintering at 350, 375, 400 and 450 °C, respectively. These values were higher than the 0.4039 nm obtained for the MA alloy. An increase in the lattice parameter and an increase in the relative intensity of the Al6Mn peaks were the result of the solid solution decomposition that occurred with the precipitation of the Al6Mn phase. However, the lattice parameter and the intensity of peaks suggest a higher content of solute elements after short-time sintering at the temperatures of 350–400 °C (Figure 9), when compared with the solute content for the granules after annealing for 2–4 h (Figure 7).
At room temperature, the brittle fracture occurred at the elastic deformation stage for the alloys processed at 350–400 °C. This meant that the ultimate strength could be determined; however, the yield strength could not be determined (Table 2). At a test temperature of 350 °C, the deformation behavior was different for the alloys sintered at 350–400 °C and 450 °C. The samples sintered at 350–400 °C did not fail during compression and showed stable flow, a yield strength was collected and is shown in Table 2. The low strength level was observed after sintering at 350 °C (Figure 10), values in the range of 160–180 MPa were observed at room temperature and 100–125 MPa at elevated temperature. The low strength can be explained by a high fraction of cavities. For the samples treated at 375 and 400 °C, the ultimate strength at room temperature increased due to annealing of the hot-pressed samples. The values of σmax were increased from 415 to 475 MPa and from 350 to 610 MPa for the alloys sintered at 375 and 400 °C, respectively. At a test temperature of 350 °C, the yield strength was also improved by heat treatment, increasing from ~160 to 190 MPa for the alloys sintered at both 375 and 400 °C.
At room temperature, a yield strength of 1100 MPa, ultimate strength of 1200 MPa and 0.5% strain at fracture were achieved for the alloy sintered in a two-step regime, hot pressed at 450 °C and further annealed at 400 °C. This alloy also had a low fraction of cavities. These samples also exhibited the highest yield strength of 380 MPa during compression at 350 °C. Meanwhile, unlike the samples sintered at lower temperatures, the samples processed at 450 °C failed after a small strain of 3.5 ± 0.5%. The mean maximum strength for these samples was 440 MPa (Table 2).
In order to explain the difference in deformation behavior at a temperature of 350 °C for the alloys treated with different regimes, the microstructure after a small compressive strain of 10% for the low-temperature sintered alloys and 3% for high-temperature sintered alloys was studied (Figure 11). For this purpose, the samples with rectangular shape and size of 5 × 5 × 7 mm were cut from the sintered samples, and one side of these samples was polished before compression. The samples sintered at 350 °C demonstrated many cavities at the boundaries of the granules, thus the strain was localized in the intergranular regions (Figure 11a). For the samples sintered at the higher temperatures of 375 and 400 °C, the surface relief also indicated cavitation development at the granular boundaries during deformation. Strain was localized in intergranular areas, but intragranular deformation with many folds on the surface was also observed in the body of the granules (Figure 11b,c). For the sample sintered at 450 °C, the strain was uniform and localized in the granular body, and only a few cavities/microcracks were observed at the granular boundaries during compression (Figure 11d).

4. Discussion

4.1. Effect of Mechanical Alloying

The main processes observed during the mechanical alloying of the studied Al–7.7Mn–4.9Zr–3.2Cu alloy are the grain refinement and dissolution of the alloying elements with further precipitation of the secondary Al6Mn phase after long-term milling. These processes were the same as those for the binary Al–10Mn alloy [69,70] and ternary Al–Mn–Cu alloys, with similar ~8%Mn and ~3.5%Cu [74] and with higher element contents of ~15.3%Mn and ~6.2%Cu [98]. Zirconium was completely dissolved in Al, but the precipitation of the Al3Zr phase was not detected during the mechanical alloying of the studied alloy. It is notable that precipitation of Al3Zr does not occur in a milling time range of 8 to 20 h for binary Al–Zr alloys [71,73,99]. The L12–Al3Zr phase is formed in Al–25at%Zr alloy after high energy ball milling of elemental Al and Zr powders [100].
The effect of stearic acid, used as a process control agent, was found to be important for microstructure formation during the mechanical alloying. According to the XRD analysis, the addition of stearic acid together with Zr resulted in the formation of ZrH2. Therefore, Mn and Cu atoms dissolve, but Zr atoms cannot dissolve in Al. The sample processed without stearic acid showed dissolution of all alloying elements, including Zr. These effects explain the differences in lattice parameters for the alloy with and without stearic acid. Stearic acid was found to facilitate the dissolution of Mn and precipitation of the Al6Mn phase during milling. The PCA-free alloy showed higher hardness, which can be attributed to the solid solution strengthening effect of Zr and finer crystallite size. For the PCA-free alloy after milling for 10 h, the average grain size was ~16 nm according to TEM, the same as the crystallite size determined by XRD. The crystallite size for PCA-containing alloy after the same milling time was ~26 nm. As the crystallite size did not change significantly during milling, it can be suggested that the increase in hardness from ~300 to ~500 HV for the PCA-bearing alloy and from 550 to 650 HV for the PCA-free alloy was due to the precipitation of the secondary Al6Mn phase during the milling process.
XRD data suggest that the maximum solubility of the alloying elements for the PCA-free alloy was achieved after mechanical alloying for ~10 h. Similar milling times are required to dissolve Mn and Cu during milling of the other complexly alloyed alloys [74,98]. Manganese dissolution shifts the Al diffraction peaks to the right and decreases the Al lattice parameter, but Zr shifts them to the left and increases the Al lattice parameter. Thus, the changes are opposite, making it difficult to estimate solute content from lattice parameter values. Assuming that all alloying elements are in the dissolved state, the theoretical lattice parameter should be 0.4033 nm. The Cu- and Zr-containing phases were not detected after milling for the PCA-free sample, but Al6Mn was incompletely dissolved, thus a minimum for the lattice parameter of 0.4039 nm is proposed. Copper and zirconium were completely dissolved and only 5.1 wt% (2.7 at%) Mn was present in the solid solution. The increase of the lattice parameter to 0.4042 nm after milling for 12.5 h suggests that the aluminum solid solution contains ~4 wt% (2.1 at%) Mn.
The lattice parameter should be ~0.4026 nm, assuming that all 3.2 wt% (1.4 at%) Cu and ~6 wt% (~3 at%) Mn are dissolved (which has been observed for Al–Mn and Al–Mn–Cu alloys [67,68,69,74]), but Zr is not dissolved. For the PCA-free alloy after milling for 5 h, where ZrH2 phase was detected, the lattice parameter was ~0.4033 nm, suggesting the dissolution of ~2.2 wt%Zr (0.7 at%).

4.2. Phase Composition of the Alloy after Annealing and Hot-Press Sintering

Due to the significant non-equilibrium conditions and high concentration of lattice defects for MA granules, recovery, recrystallization and grain growth can occur during annealing [101,102]. The high concentration of vacancies and density of dislocations and grain boundaries accumulated during the milling [51] should facilitate diffusion controlled processes such as precipitation of the secondary phases. Annealing reduces the dislocation density and leads to decomposition of the solid solution.
The polythermal section calculated by ThermoCalc (database TTAl5) for the studied alloy is shown in Figure 12. The chemical composition of the alloy indicates the existence of three equilibrium phases Al6Mn, Al20Cu2Mn3, and Al3Zr (D023) in the studied temperature range. The formation of the Al3Zr (L12) and CuAl2 phases is also reasonable at the temperatures studied [35,103,104].
The Al6Mn, CuAl2, Al3Zr (L12), and Al20Cu2Mn3 phases were revealed by XRD after annealing in the studied temperature range of 350–450 °C for 2–4 h. During hot pressing, the maximum time at elevated temperature was ~15 min. At elevated temperatures of 400–450 °C, the Al6Mn phase precipitated, a minor amount of the Al3Zr and Al20Cu2Mn3 phases were found; however, the other secondary phases were not defined by XRD. At the same time, when the precipitates are fine and the fraction is small, it may be difficult to identify phases by XRD, so other phases are possible.
Thermal analysis revealed several thermal effects in mechanically alloyed granules upon heating. The exothermic P1 peak observed at 140–250 °C was probably the result of the recovery/recrystallization processes [101,102]. The low-temperature effect can also be the result of the θ-phase (CuAl2) precipitation that was observed by XRD for the alloy annealed at 350–450 °C for 2–4 h. The P1 peak was non-symmetric with two maxima, so both phenomena were possible. The P2 effect was the result of Mn-bearing phases that can precipitate at temperatures of 300–350 °C [103,104,105,106,107]. The P3 effect with the maximum at ~430 °C can be explained by the precipitation of the Al3Zr phase, which is in good agreement with several studies showing that L12–Al3Zr phase precipitates predominantly at 350–450 °C [37]. The P4 peak observed near 500 °C can be related to the formation of the stable D023–Al3Zr phase [37].

4.3. Mechanical Properties of the Alloy after Annealing and Hot-Press Sintering

Analysis of hardness evolution during annealing at the temperatures of 350–450 °C for mechanically alloyed granules showed that competitive processes occurred in the microstructure. Softening after annealing for 0.5–1 h was the result of a decrease in the level of lattice defects and grain growth. At higher temperatures, the decrease was more pronounced, which may be the result of the recrystallization and intensive grain growth. Further hardening is explained by the formation of fine precipitates. The slight increase in hardness observed at 450 °C after annealing for 2–4 h can be explained by the coarsening of precipitates and more intensive grain growth at elevated temperature.
The density of the alloys increased from ~2.7 to ~3.1 g/sm3 and the fraction of cavities decreased from ~12 to ~1.5%, with an increase in sintering temperature from 350 to 450 °C. The theoretical density can be estimated in two ways. Given that all components are not dissolved in the matrix and exist as a mixture of pure metals, the density can be estimated additively. For the composition studied, the density of the alloy should be ~3.4 g/cm3. This value suggests about 6% of cavities after sintering at 450 °C, which is higher than the experimental values of ~2%. Considering solute elements after milling for 10 h (~5%Mn, 4.9%Zr, 3.2%Cu), the density of Al solid solution can be calculated from the XRD data. The lattice parameter was 0.4039 nm after milling for 10 h. Therefore, the density ρ can be estimated by Equation (1).
ρ = Ass ∙ n/(Vss ∙ N)
where Ass is the mean atomic weight of the solid solution, n is the number of atoms per unit cell, Vss is the volume of the unit cell, and N is Avogadro’s constant.
The Ass is calculated additively using the concentrations and atomic weight of the elements in the studied alloy. The n = 4 is taken for the f.c.c. lattice and the Vss has been estimated using a lattice parameter. According to Equation (1) the density of the solid solution is ~3.04 g/cm3. The density of the remaining Al6Mn phase with a volume fraction of ~9% (for undissolved 2.7%Mn) is 3.32 g/sm3; thus, the theoretical density of the alloy is 3.06 g/sm3, which is similar to the maximum experimentally measured values after sintering at 450 °C.
At room temperature, a high fraction of cavities was the cause of brittle fracture and low strength compared with the high microhardness values of the low-temperature sintered alloys. Microhardness is less sensitive to the cavities, but these macrodefects initiate cracks and lead to their progressive growth. A high fraction of cavities and brittle fracture resulted in a low strength for the low-temperature sintered alloys. The hardness level of ~500 HV for low-temperature sintering suggested a yield strength of ~1600 MPa (9.8 × HV/ σ Y S = 3 [108]), while a high fraction of the cavities decreased the strength level for the low-temperature sintered alloys. This demonstrates that the sintering temperature should be above 400 °C to increase the density of the bulk samples but higher temperatures have a negative effect on the strength level. The maximum yield strength of ~1100 MPa was observed for the alloys sintered at 450 °C when the fraction of cavities was small.
At a compression temperature of 350 °C, the deformation behavior was different for the samples sintered at 350–400 °C and 450 °C. A high fraction of cavities (free volume) in the low-temperature sintered samples resulted in intergranular deformation and apparently high ductility. For the sample sintered at the elevated temperature with a low fraction of cavities (small free volume), the deformation behavior changed. Strain was localized in the granules, strain hardening occurred, cracks formed, and the samples failed. Thus, the effect of the apparently high ductility of the samples processed at low temperatures was the result of the high fraction of cavities and the intergranular deformation accommodated by cavitation. The yield strength values of these samples do not primarily characterize the strength of the alloy, but rather the strength of the intergranular bonds that have formed during sintering. Importantly, annealing at a sintering temperature increased density of the hot-pressed samples, which improved the strength properties. The reason for this phenomenon is an increased intergranular bonding force due to the rapid diffusivity in the nanostructured alloy. The deformation behavior was controlled by intergranular sliding, a process that is similar to grain boundary sliding in superplastic flow. For the studied materials, deformation is controlled by sliding at weak intergranular boundaries between high-strength MA granules. Higher sintering temperatures result in a lower fraction of cavities and deformation occurring in the body of the granules. The surface relief with multiple fine folders and the strain hardening effect in the stress–strain curves suggest dislocation-induced mechanisms of deformation.
The room-temperature strength of the studied alloy is significantly higher than for Al–Cu/Al–Cu–Mn-based alloys and similar to high-strength Al–Fe powder alloys and Al–matrix composites with nanoparticles (Table 3).
For low-alloyed Al–Mn–Cu–Zr(Cr,Sc) alloys (4–7 wt% of the alloying elements) strengthened by Al20Cu2Mn3, icosahedral quasicrystalline and L12 precipitates, an ultimate tensile strength of ~600 MPa at room temperature [20,110] and a yield strength of 150 MPa at elevated temperature of 350 °C [105] are reported. Compressive strengths of 800 MPa at room temperature and 350 MPa at 300 °C are observed for hot-press sintered MA granules of Al–5at/10wt% Zr alloy [73]. Notably, the Zr-free Al–7.7Mn–3.1Cu (wt%) alloy processed by high-energy ball milling and hot-press sintering at 400 °C exhibits a yield strength of ~900 MPa at room temperature and ~220 MPa at 350 °C [109]. Similar strength are found for aluminum-based alloys mechanically alloyed with Al3Zr and Al3Ti particles [100]. The studied Al–4.2 at/7.7wt%Mn–1.5 at/3.2 wt%Cu–1.6 at/4.9 wt%Zr alloy demonstrated a yield strength of 1100 MPa, an ultimate compressive strength of 1200 MPa at room temperature, and a yield strength of 380 MPa and an ultimate compressive strength of 440 MPa at the elevated temperature of 350 °C. Therefore, Zr has a considerable influence on the strength properties of the alloy at both temperatures. Similarly high compressive strengths of 1000–1200 MPa at room temperature and 400–500 MPa at 350 °C are observed for the Al–5 at/10 wt%Fe alloy processed by mechanical alloying and SPS at 480 °C [66]. The high strength of the Al–Fe alloy is provided by its nanocrystalline structure, fine Al6Fe intermetallic phase, and the oxide and carbide precipitates introduced by PCA [66]. A high content of Mn, Zr, precipitates with these elements and the nanocrystalline structure provided by mechanical alloying provided a high strength level for the studied Al–Mn–Cu–Zr PCA-free alloy.
The alloy studied has a high strength, but its low ductility may limit the fatigue life [113,114]. The alloy has the structure of a natural composite reinforced by fine intermetallic precipitates and does not contain coarse ceramic particles that facilitate cracking, suggesting an increased fatigue life compared with ceramic-reinforced composites. Fine precipitates may also improve crack resistance [115]. Meanwhile, further experiments are required to investigate the fatigue properties of the studied alloy.

5. Summary

The novel Al–7.7Mn–3.2Cu–4.9Zr (wt%) alloy was processed by mechanical alloying (MA) with and without stearic acid as a process control agent (PCA) followed by hot-press sintering at temperatures in the range of 350–450 °C. The microstructure, phase composition, and mechanical properties of mechanically alloyed granules and sintered bulk samples of the PCA-free alloy were studied. The main conclusions are as follows.
  • The dissolution of Mn, Zr, and Cu with further precipitation of the Al6Mn phase occurred during mechanical alloying. For the PCA-free alloy, the dissolution of Zr, Cu, and the Mn solute content of about 5 wt% were observed after milling for 10 h. The grain size measured by TEM was 16 nm, which was the same as the crystallite size estimated by XRD. Due to the high solute content, fine precipitates and nano-grained structure, the MA granules exhibit a hardness of ~500–650 HV.
  • The addition of stearic acid helped to refine the granules and facilitated Mn dissolution and precipitation of the secondary Al6Mn phase during milling but resulted in the formation of zirconium hydrate ZrH2 at early milling stages. Stearic acid decreased Zr solubility, increased the crystallite size of the MA granules to ~22 nm and decreased hardness to ~300–500 HV.
  • According to XRD data, annealing of the MA granules at the temperatures of 350–450 °C for 2–4 h resulted in decomposition of Al solid solution with precipitation of the Al6Mn, L12–Al3Zr, Al20Cu2Mn3, and Al2Cu secondary phases. The crystallite size increased to 26–33 nm and hardness decreased from 530 HV for as-milled condition to 380–500 HV after annealing.
  • Sintering and further annealing at temperatures of 350–400 °C provided a high hardness of 500–520 HV but a low density of ~2.7–2.9 g/sm3 and a large fraction of cavities of 6–12%. As a result, the studied alloy has a brittle fracture and low strength values at compression tests. In the elevated temperature test at 350 °C, the alloy exhibited an apparently stable flow with a low yield stress of 100–190 MPa and without the failure associated with the intragranular deformation mechanism accommodated by cavitation.
  • Hot-press sintering at 450 °C, followed by annealing at 400 °C for 2 h resulted in a high density of ~3.1 g/sm3, which was similar to theoretical values, and a low fraction of cavities of ~1.5%. This sintering regime provided a lower hardness of 400 HV but high strength properties of the alloy in compression tests; a yield strength of 1100 MPa, an ultimate strength of 1200 MPa, and a strain at fracture of 0.5% were observed at room temperature, and a yield strength of 380 MPa, an ultimate strength of 440 MPa, and a strain at fracture of 3.5% were observed at 350 °C. The SEM studies of the surface structure after small strains at 350 °C showed that the deformation for the high-temperature sintered alloy was localized in the granular body.

Author Contributions

Conceptualization, A.V.M. and A.S.P.; methodology, O.A.Y., A.S.P., A.I.B. and N.B.E.; software, A.I.B. and N.B.E.; formal analysis, A.V.M. and N.B.E.; investigation, O.A.Y., A.I.B., A.S.P. and N.B.E.; data curation, N.B.E. and O.A.Y.; validation, A.G.M., N.B.E. and A.V.M.; funding acquisition, A.V.M.; resources O.A.Y. and N.B.E.; supervision, A.V.M.; writing—original draft preparation, O.A.Y., A.G.M. and A.V.M.; writing—review and editing, A.V.M.; project administration, O.A.Y. and A.V.M. All authors have read and agreed to the published version of the manuscript.

Funding

This research was funded by the Russian Science Foundation (RSF), grant number 23-19-00791.

Data Availability Statement

The raw data supporting the conclusions of this article will be made available by the authors on request.

Acknowledgments

The authors wish to thank the MISIS Collective Use equipment Center “Material Science and Metallurgy” and Nataliya Tabachkova for TEM studies.

Conflicts of Interest

The authors declare no conflict of interest.

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Figure 1. XRD patterns for the (a) alloy with PCA and (b) without PCA (wt.%) alloy after high-energy ball milling for different times. (c) The magnified diffraction peak for (420)Al at different treatments.
Figure 1. XRD patterns for the (a) alloy with PCA and (b) without PCA (wt.%) alloy after high-energy ball milling for different times. (c) The magnified diffraction peak for (420)Al at different treatments.
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Figure 2. The microstructure of the alloy (a,b,e) without and (c,d) with PCA after milling for (a,c,e) 5 h and (b,d,f) 10 h (secondary electron images) with (e,f) EDS element distribution maps for the marked areas (backscatter electron images).
Figure 2. The microstructure of the alloy (a,b,e) without and (c,d) with PCA after milling for (a,c,e) 5 h and (b,d,f) 10 h (secondary electron images) with (e,f) EDS element distribution maps for the marked areas (backscatter electron images).
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Figure 3. The evolution of the (a) lattice parameter, a; (b) crystallite size, CSR; and (c) microhardness, HV0.025, during mechanical alloying for the alloys with PCA (red-colored lines) and without PCA (black-colored lines).
Figure 3. The evolution of the (a) lattice parameter, a; (b) crystallite size, CSR; and (c) microhardness, HV0.025, during mechanical alloying for the alloys with PCA (red-colored lines) and without PCA (black-colored lines).
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Figure 4. TEM images of the granules after mechanical alloying for 10 h. (a) Bright field, (b) dark field images and (c) SAED for the same area.
Figure 4. TEM images of the granules after mechanical alloying for 10 h. (a) Bright field, (b) dark field images and (c) SAED for the same area.
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Figure 5. The DTA curves of the studied alloy after 10 h of milling.
Figure 5. The DTA curves of the studied alloy after 10 h of milling.
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Figure 6. Microhardness evolution of the PCA-free alloy that had been milled for 10 h after annealing at the temperatures of 250–450 °C.
Figure 6. Microhardness evolution of the PCA-free alloy that had been milled for 10 h after annealing at the temperatures of 250–450 °C.
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Figure 7. (a) XRD data evolution after annealing of the MA PCA-free alloy. (b) The magnified diffraction peak for (420)Al after different treatments.
Figure 7. (a) XRD data evolution after annealing of the MA PCA-free alloy. (b) The magnified diffraction peak for (420)Al after different treatments.
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Figure 8. SEM secondary electron images of the hot-pressed samples at (a) 350 °C, (b) 375 °C, (c) 400 °C, and (d) 450 °C and further annealing for 2 h.
Figure 8. SEM secondary electron images of the hot-pressed samples at (a) 350 °C, (b) 375 °C, (c) 400 °C, and (d) 450 °C and further annealing for 2 h.
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Figure 9. (a) XRD patterns for the samples hot pressed at temperatures of 350 °C, 375 °C, 400 °C, and 450 °C. (b) The magnified diffraction peak for (420)Al at different treatments.
Figure 9. (a) XRD patterns for the samples hot pressed at temperatures of 350 °C, 375 °C, 400 °C, and 450 °C. (b) The magnified diffraction peak for (420)Al at different treatments.
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Figure 10. The stress–strain curves at 350 °C for the samples hot-press sintered and further annealed at 350–450 °C.
Figure 10. The stress–strain curves at 350 °C for the samples hot-press sintered and further annealed at 350–450 °C.
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Figure 11. The surface of the pre-polished samples of the PCA-free alloy sintered at (a) 350, (b) 375, (c) 400 and (d) 450 °C after compression for (a–c) 10% and (d) 3% at a temperature of 350 °C.
Figure 11. The surface of the pre-polished samples of the PCA-free alloy sintered at (a) 350, (b) 375, (c) 400 and (d) 450 °C after compression for (a–c) 10% and (d) 3% at a temperature of 350 °C.
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Figure 12. The polythermal section of the Al–4.9Zr–3.2Cu–xMn equilibrium phase diagram, the red dashed line shows the studied alloy’s composition (ThermoCalc).
Figure 12. The polythermal section of the Al–4.9Zr–3.2Cu–xMn equilibrium phase diagram, the red dashed line shows the studied alloy’s composition (ThermoCalc).
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Table 1. Lattice parameter, CSR in granules of the studied PCA-free alloy after mechanical alloying for 10 h, annealing of the MA granules and hot-press sintering.
Table 1. Lattice parameter, CSR in granules of the studied PCA-free alloy after mechanical alloying for 10 h, annealing of the MA granules and hot-press sintering.
TreatmentLattice Parameter a, nmCoherent Scattering Region, CSR, nm
Mechanical alloying0.403917 ± 3
Annealing at 350 °C for 2 h0.404823 ± 5
Annealing at 350 °C for 4 h0.405531 ± 6
Annealing at 400 °C for 2 h0.405526 ± 5
Annealing at 450 °C for 2 h0.405329 ± 8
Hot pressing at 350 °C 0.404234 ± 2
Hot pressing at 375 °C0.404325 ± 5
Hot pressing at 400 °C0.404733 ± 7
Hot pressing at 450 °C0.405333 ± 8
Table 2. Properties for the sintered samples of the studied PCA-free alloy.
Table 2. Properties for the sintered samples of the studied PCA-free alloy.
Consolidation Temperature and
Heat Treatment Regime		HV, kgs/mm2Density, g/sm3Ultimate Compressive Strength (σmax) at 25 °C, MPaYield Strength (σ0.2) at 350 °C, MPa
350 °C500 ± 302.70 ± 0.03180 ± 10125 ± 20
350 °C and annealing at 350 °C for 2 h475 ± 302.69 ± 0.02160 ± 20100 ± 15
375 °C520 ± 202.77 ± 0.02415 ± 15160 ± 15
375 °C and annealing at 375 °C for 2 h515 ± 202.89 ± 0.02475 ± 15190 ± 10
400 °C470 ± 202.78 ± 0.02350 ± 30165 ± 10
400 °C and annealing at 400 °C for 2 h510 ± 152.85 ± 0.01610 ± 25185 ± 15
450 °C380 ± 203.05 ± 0.02--
450 °C and annealing at 400 °C for 2 h400 ± 153.15 ± 0.051100 ± 20 */1200 ± 30 **380 ± 12 */440 ± 13 **
* The yield strength value; ** The ultimate compression strength value.
Table 3. Mechanical properties of several aluminum-based alloys and composite materials at room and elevated temperatures (temperature is noticed in the index).
Table 3. Mechanical properties of several aluminum-based alloys and composite materials at room and elevated temperatures (temperature is noticed in the index).
Alloy/CompositeHardness, HV, kgs/mm2Yield Strength (RT/Elevated
Temperature), MPa		Ultimate Compressive or Tensile Strength, (RT/Elevated
Temperature) MPa 		Strain–to–Fracture, (RT/Elevated
Temperature)%		Ref.
Al–7.7 wt%Mn–3.2 wt%Cu–4.9 wt%Zr530–6301100RT/3803501200RT/4403500.5RT/3.5350This study
Al–7.7 wt%Mn–3.1 wt%Cu500870RT/220350900RT/350350 0.27RT/40350[109]
Al–3.17%Mn–3.08 %Cu–0.35%Be–0.21%Sc–0.14%Zr~110304RT/150350437RT/155350180350[105]
Al–Mn–Cu–Zr-450RT590RT0.6RT[20]
Al–3%Mn–1%Cr–2%Cu-≈300RT600RT≥20 RT[110]
Al–10 wt%Zr270310300822RT/3443000.1RT/0.2300[73]
Al–10 wt%Fe–PCA-1000RT/4003501200RT/5003500.1–0.2RT/≥15350[66]
Al alloy + graphene NP--1200RT/110350-[111]
7075 + Ag–C NPs300---[12]
2024 + 10%Al2O3-460RT/325180495RT/3401802.0RT/6.6180[112]
Al–5 wt%Zr–10%Al203520 220350340350 25350[72]
AA6063 + 1 wt%Cu + Al3Ti265872RT/273300~970RT/290300~8300[100]
AA6063 + 1 wt%Cu + Al3Zr265841RT/242300~900RT/254300~5300[100]
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Yakovtseva, O.A.; Mochugovskiy, A.G.; Prosviryakov, A.S.; Bazlov, A.I.; Emelina, N.B.; Mikhaylovskaya, A.V. The Microstructure and Properties of Al–Mn–Cu–Zr Alloy after High-Energy Ball Milling and Hot-Press Sintering. Metals 2024, 14, 310. https://doi.org/10.3390/met14030310

AMA Style

Yakovtseva OA, Mochugovskiy AG, Prosviryakov AS, Bazlov AI, Emelina NB, Mikhaylovskaya AV. The Microstructure and Properties of Al–Mn–Cu–Zr Alloy after High-Energy Ball Milling and Hot-Press Sintering. Metals. 2024; 14(3):310. https://doi.org/10.3390/met14030310

Chicago/Turabian Style

Yakovtseva, Olga A., Andrey G. Mochugovskiy, Alexey S. Prosviryakov, Andrey I. Bazlov, Nadezhda B. Emelina, and Anastasia V. Mikhaylovskaya. 2024. "The Microstructure and Properties of Al–Mn–Cu–Zr Alloy after High-Energy Ball Milling and Hot-Press Sintering" Metals 14, no. 3: 310. https://doi.org/10.3390/met14030310

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