Materials with nanosize- or ultrafine-grained structure have received much attention in the last decades. Their yield stress is increased according to the Hall-Petch relationship. Their strength characteristics can be further improved by the introduction of a dense and homogeneous distribution of fine precipitates or dispersoids, which not only contribute to the materials’ strength, but also stabilize the fine-grained microstructure. These particles can successfully pin grain boundaries and dislocations, and suppress recrystallization and grain growth.
Materials with a fine-grained structure can be, among others, prepared by a powder metallurgy (PM) route. Gas atomization based on the disintegration of a molten metal stream by high pressure gas jets is a typical representative of rapid solidification methods which is capable of producing fine powder particles. High cooling rates (103
K/s) along with a high surface-to-volume ratio of droplets lead to a fine microstructure, extended solid solubility, and to the formation of metastable phases [1
], all enhancing material strength. Further microstructural refinement can be achieved by the mechanical milling of atomized powders. Repeated flattening, fracturing, and rewelding of powder particles occurs during the milling process, and a large amount of deformation energy is introduced into the powder material. The very high strain rates lead to severe plastic deformation and high dislocation density. The self-organization of dislocations into cell networks, low-angle boundaries, and finally high-angle boundaries is considered to be the mechanism of grain refinement down to the nanoscale range [2
]. Additionally, fine oxide and nitride particles from the powder particle surface, atmosphere, or milling agent can be introduced into the matrix during the milling process, which can further enhance strength through dispersion strengthening [5
]. These dispersoids can also play an important role in maintaining powders’ fine-grained structure during a subsequent consolidation process.
The microstructure and properties of PM materials can be strongly affected by the selected consolidation process. The most frequently used sintering methods, such as hot pressing (HP) or hot isostatic pressing (HIP), are based on concurrent external heating and pressing (e.g., [6
]). The elevated temperature and prolonged heat exposure characteristic of these sintering methods can lead to recrystallization, grain growth, or massive precipitation, all decreasing the gains from atomization and milling. Spark plasma sintering (SPS) represents an alternative method that can minimize these detrimental effects. SPS combines uniaxial pressure with heating by a low-voltage direct current (DC) flowing through the sample (in the case of conductive material). A large Joule heat is evolved especially at the contact points between powder particles, and the temperature at these points can highly exceed the set one. It is claimed that SPS affects particularly the surfaces of the powder particles, and the particle cores can retain their original microstructure [8
]. During SPS, the oxide layers, which are regularly present at the surfaces of powder particles, can be disrupted more efficiently than by conventional sintering techniques [7
]. SPS is thus able to produce fully dense compacts at lower temperatures and in substantially shorter times than other consolidation routes, and the undesirable processes, such as recrystallization or grain growth, can be restricted.
PM processing has been successfully used in many aluminium-based materials. Both gas atomized (e.g., [11
]) and milled (e.g., [16
]) powders have been sintered by a variety of methods. SPS results in the lowest porosity [20
], and is also very efficient in the production of a very fine-grained microstructure [19
]. Numerous experiments have been performed on the 7xxx series of Al-based alloys. A combination of gas atomization and SPS has been used for the preparation of the Al7075 alloy with the relative density exceeding 99%. The grain size of this material was in the range of between 1 and 10 µm, and remained stable even after annealing at 500 °C [21
]. Even finer Al7075 material with a grain size at the nano-meter scale has been prepared by a combination of milling of elemental powders followed by hot pressing [18
Zirconium is frequently added to the Al7075 alloy for an improved stabilization of the fine-grained structure. Small precipitates of the metastable Al3
Zr phase serve as grain boundary pinners [23
]. A grain size slightly exceeding 1 µm has been observed in the Al7075 + 0.7 wt % Zr [13
] and 1 wt % Zr alloys [14
] prepared by a combination of gas atomization and hot extrusion, and both these materials exhibited high strain rate superplasticity. Furthermore, the introduction of Al3
Zr particles displaced the border of thermal stability to higher temperatures in a cryorolled 7075 alloy in [24
]. A nano-sized microstructure has been observed also in the Al7068 + 0.1 wt % Zr alloy consolidated by hot pressing [17
]. Unfortunately, the relative density of this material was below 97%, which resulted in a brittle behaviour.
Whereas the commercial Al alloys are doped usually only by a few tenths wt % of Zr, the Al7075 alloy used in our research is a model material where a higher Zr content was chosen in order to emphasize the role of Zr addition in microstructural stabilization. A combination of gas atomization, mechanical milling of the powders, and spark plasma sintering was expected to result in a very fine-grained material with minimum porosity, good microstructural stability, and good mechanical properties. A proper investigation of the microstructure and phase composition after each step of processing was performed using a variety of experimental techniques, and the findings were correlated with the microhardness.
2. Materials and Methods
The Al7075 alloy powder modified by the addition of about 1 wt % of Zr was prepared by Nanoval GmbH and Co. KG, Berlin, Germany. The alloy was atomized by nitrogen and sieved down to 50 µm. The average particle size was stated by the producer to be 20.9 µm. The atomized powder will be further denoted as Al7075Zr_A. The chemical composition of the gas-atomized powder according to X-ray fluorescence (XRF) spectroscopy is given in Table 1
High energy ball milling of the atomized powder was carried out in a Retsch PM 100 CM mill (Retsch GmbH, Haan, Germany) at room temperature. The atomized powder was dry milled in a steel tank with a speed of 400 revolutions per minute (RPM) for 8 h in an Ar atmosphere using 20 mm steel balls. The ball-to-powder ratio was 40:1. The milled powder is further denoted as Al7075Zr_M.
Both the atomized and milled powders were compacted by SPS using FCT SPS-HP25 (FCT Systeme GmbH, Effelder-Rauenstein, Germany). The scheme of the sintering process is presented in Figure 1
. During consolidation, the samples were free heated up to 400 °C, and the 425 °C sintering temperature was reached with a heating rate of 25 °C/min. Simultaneously with temperature, the uniaxial pressure was increased up to 80 MPa. After the holding time of 4 min, the compact was free cooled and unloaded. The density of SPS compacts was measured using He pycnometry, and the sample porosity was additionally controlled by image analysis of SEM micrographs (10 images were evaluated for each material). The compacts prepared from the atomized and milled powders are further referred to as Al7075Zr_AC and Al7075Zr_MC, respectively.
The phase composition and structural changes were investigated by X-ray diffraction (XRD) using a diffractometer D8 Discover (Bruker AXS, Brno, Czech Republic) in Bragg-Brentano geometry with a CuKα source and a NiKβ radiation filter. A quantitative Rietveld analysis was performed by TOPAS V5 (Bruker AXS) to determine the weight fraction of all identified phases. The sizes of coherently diffracting domains (CDD) and microstrains were evaluated from the broadening of diffraction peaks assuming that small crystallites and microstrains contribute to the broadening of the Lorentzian and Gaussian components of the pseudo-Voigt function, respectively [25
The morphology and microstructure of the powders were investigated by light microscopy (LM) using an Olympus GX51 (Olympus, Prague, Czech Republic) microscope. The powders were cold mounted in acrylic cold mounting resin, metallographically processed, and finally etched by Dix-Keller reagent to reveal microstructure details.
Finer microstructural details were investigated by scanning electron microscopy (SEM) using an FEI Quanta 200F microscope (FEI, Brno, Czech Republic) equipped with a field emission cathode (FEG). Chemical composition was studied by energy dispersive X-ray spectroscopy (EDS) (EDAX, Nachod, Czech Republic). The grain size and grain orientation were analysed using electron backscattered diffraction (EBSD). The EBSD camera DigiView3 (EDAX, Nachod, Czech Republic) was used for signal detection, and TSL OIM 7 software (Version 7.3.0, EDAX, Nachod, Czech Republic) was used for data analysis. The preparation of samples for SEM included hot mounting (180 °C, 2.5 min) in conductive resin with carbon filler (in the case of powder samples), mechanical grinding, mechanical polishing up to 1 µm, chemical polishing with SiO2 oxide in polishing suspension, and electro-polishing (just for EBSD).
Transmission electron microscopy (TEM) investigations were performed using a JEOL 2200FS microscope (JEOL, Peabody, USA) operated at an accelerating voltage of 200 kV. TEM specimens from powders were produced using focused ion beam (FIB) milling in a Zeiss Auriga Compact SEM microscope. The samples were extracted from a chosen powder particle embedded in conductive resin with carbon filler; further, they were milled to electron transparency and mounted onto a copper TEM grid. Interplanar spacings and angles were measured on electron diffraction patterns taken along reasonably low index axis orientations. These were then matched to crystal structures selected from inorganic crystal structure database (ICSD) [26
] based on the composition criteria (EDX). Finally, the whole pattern was simulated in JEMS [27
] and a match with the experimental pattern was verified.
From the SPS compacts, discs with a diameter of 3 mm were cut from a plane parallel to the direction of load applied during SPS. Thin foils for TEM observations were prepared by electrochemical thinning using a double-jet polisher, Tenupol-5, under 15 V at −15 °C in a solution of 33% nitric acid in methyl alcohol.
The strength of the materials was characterized using microhardness measurements. An automatic microhardness tester Qness Q10A+ with a Vickers indentor (Qness GmbH, Golling, Austria) was used. For the powder materials, one indent with a load of 10 N was applied to an individual particle embedded in the cold mounting resin, and the mean microhardness value was calculated from measurements performed on different powder particles. For the compact materials, indentations were performed with the applied load of 50 N on planes parallel to the direction of the applied stress during SPS. In order to obtain good statistics, areas of 4.4 × 4.4 mm2 were investigated, with a distance between individual indents of 200 µm.
The Al7075 alloy is a commercial material, which is frequently used also for research purposes. This alloy originates its high strength especially from the precipitation sequence occurring either at room temperature or at elevated temperatures after solution treatment. An additional strength increase can be achieved by strain hardening, i.e., by the application of a large plastic deformation. Last but not least, the reduction of the grain size can contribute to a strength increase. Whereas the former treatments can be applied to materials prepared through traditional processing routes, substantial grain refinement requires a special attitude. The minimum grain size achievable by a thermomechanical treatment applied to an ingot cast material is above 10 µm [28
]. The application of severe plastic deformation by equal channel angular pressing (ECAP) made it possible to reduce the grain size down to approximately 1 µm; however, the fraction of high-angle boundaries is usually only slightly above 50%, so that many of these “grains” should be more correctly described as subgrains [30
]. Powder metallurgy represents an alternative route in the processing of very fine-grained materials. Both rapid solidification methods and the mechanical milling of powders were used in our research.
Gas atomization is a typical representative of a rapid solidification technique. A high temperature gradient in the liquid phase and a high rate of solidification front movement are typical for the onset of solidification, and can result in a featureless (segregation-free) microstructure in very small droplets. In our material, such a segregation-free microstructure was also found in satellites attached to some coarser droplets (see insert in Figure 2
a and Figure 3
a). The connection of smaller solid powder particles onto larger melt droplets can be attributed to a turbulent flow of particles and gas during atomization. A gradual decrease in the temperature gradient and in the rate of the solidification front results in a breakdown of planar solidification front and a formation of cellular, columnar, or dendritic microstructures [31
]. A mixture of dendritic and cellular microstructure has been observed in the gas-atomized Al7075 alloy, with a mean droplet size of about 50 µm [21
]. Due to the smaller mean droplet size of our material (about 20 µm), the cellular microstructure is the prevailing microstructure type. The cell size in the droplets is in the order of µm (Figure 3
b), which is comparable to the grain size of the materials prepared using ECAP from ingot cast material. Nevertheless, most of the cells are separated by high-angle grain boundaries.
The powder’s microstructure was completely changed by mechanical milling. The milling process consists of the flattening, fracturing, and rewelding of powder particles, which leads to alterations in powder particles’ morphology, size, and microstructure. The former spherical powder particles with a size below 20 µm became irregular-shaped, and their size increased to up to 200 µm during the milling process. Contrary to the gas-atomized powder, it was not possible to obtain diffraction patterns from the milled powder during EBSD measurements. The explanation can be found in a too-small size of the coherently diffracting volume. According to [35
], the minimum spatial resolution for aluminium is approximately 60 nm using tungsten-filament SEM and 20 to 30 nm using FEG-SEM. Both the evaluation of the crystallite size from the broadening of XRD peaks and the TEM investigation of the milled powder particles confirmed a nanocrystalline microstructure with the grain size between 20 and 100 nm (Figure 8
), which is comparable with the spatial resolution limit for EBSD. A similar grain size was observed in an AlZnMgCuZr alloy after mechanical milling for 40 h from elemental powders (27 nm) [17
] in an Al7075 alloy after cryomilling for 10 h (28 nm) [36
] and in the same alloy milled at room temperature for 20 h (26 nm) [18
]. A comparison with our results shows that the rate of grain refinement is significantly higher in our case.
Spark plasma sintering was used for the consolidation of powders. The nature of SPS is not fully understood, but a range of models can be found in the literature [37
]. It is believed that, beside the applied pressure, local electric and magnetic fields can activate the powder’s surface, help to break up and remove oxide layers present on the surface of powder particles handled without an inert atmosphere, and promote full densification. Our results confirmed (in accordance with previous experiments on other Al-based alloys [20
]) that SPS is very efficient in the preparation of compacts with nearly full density. The slightly lower density of the Al7075Zr_AC compact (as compared to the theoretical value) may be caused by the presence of some retained oxide particles rather than by porosity. The slightly higher density of the Al7075Zr_MC compact reflects probably the contamination by Ni and Ti during milling. Nearly negligible porosity was verified also by an image analysis of the SEM micrographs.
As already mentioned in the introduction, the temperature gradient along the powder particle radius causes necks to be formed between adjacent powder particles leading to material densification, whereas the powder particle interiors are not significantly influenced and retain their original microstructure. This has been observed in many materials, including a gas-atomized Al7075 alloy [21
]. The present experiments revealed that no massive grain growth occurred during SPS of the gas-atomized powder (Figure 12
a). The smaller grains observed in the Al7075Zr_AC compact correspond to cells present in the original coarser gas-atomized droplets; the size of the coarser grains in the Al7075Zr_AC compact corresponds to the size of the original smaller droplets present in the atomized material. It can be expected that these small droplets had originally a featureless microstructure due to a higher rate of solidification. Although two types of interfaces are present in the compact (original cell boundaries and boundaries between original droplets), both of them predominantly have a high-angle character (Figure 14
Mechanical milling introduces a severe plastic deformation into powder particles. The stored deformation energy represents a strong driving force for recrystallization, which can be expected to occur during SPS at 425 °C (despite a very short exposition to elevated temperatures). All experimental methods used in our research proved that there were significant microstructure changes in the Al7075Zr_MC compact. TEM revealed that the microstructure consisted of recrystallized equiaxed grains with a size on the order of 100 nm. XRD documented an increase in the crystallite size and a decrease in strain in comparison with the milled powder. Finally, the EBSD experiments showed a very homogeneous sub-microcrystalline microstructure. Grains with a mean size of below 400 nm are separated predominantly by high-angle grain boundaries. This structure is significantly finer and more homogeneous than that observed in the Al7075Zr_AC compact (Figure 12
a) or in the ECAP material [30
The mechanical properties of the Al7075-based alloys are strongly influenced by their phase composition, which is rather complicated. It has been found in ingot metallurgical material that the main alloying elements Zn and Mg are engaged in the following precipitation sequence: supersaturated solid solution-Guinier-Preston zones-metastable η’phase-stable η-MgZn2
]. The solution treatment is applied at temperatures above 470 °C, which is slightly above the SPS sintering temperature used in our experiment. The maximum strength is observed in the material containing Guinier-Preston zones and metastable η’particles. Such phase composition can be formed either during annealing at temperatures close to 140 °C or during a long-term stay at room temperature, which was the case for our materials. However, this precipitation sequence can be modified by chemical composition, the presence of microstructure defects, or further sites for heterogeneous nucleation (e.g., primary particles rich in Fe or Si), and therefore also other complex phases containing e.g., Cu can be formed. Additionally, our material is doped with Zr, which should form Al3
Zr particles stabilizing the fine-grained microstructure [23
]. It can be expected that precipitation sequences are strongly influenced by non-traditional processing routes, such as rapid solidification, milling, or sintering at elevated temperatures. Therefore, the phase composition after different processing steps was investigated.
The EDS experiments performed on the gas-atomized powder revealed the presence of Zn, Mg, and Cu in the cell boundaries of most powder particles. The peaks of phases of the type Mg(Zn,Cu,Al)2
were observed in XRD. Both phases probably also contain Cu; however, this hypothesis cannot be proved by XRD because of the very low weight fraction of both these phases (bellow 1%). Similar results have been found in the gas-atomized Al7075 alloy with a coarser droplet size [21
]. No peaks of the Al3
Zr phase were observed despite a nearly negligible equilibrium solubility limit of Zr in Al [45
]. It can be supposed that the Zr was nearly completely dissolved in the Al-based matrix due to rapid solidification. A similar result has been found in a melt-spun AlZnMgCuZr alloy [46
The XRD experiment revealed no peaks of any intermetallic phases in the as-milled powder. The precipitates present in the gas-atomized powder were dissolved or their fraction decreased to under the limit of detectability of the used method. Simultaneously, as a result of the dissolution of the alloying elements in the Al-matrix, the Al peak was slightly shifted. The dissolution of the alloying elements altered the lattice parameters of the matrix. In accordance with the element's atomic radius, Zn and Cu shifted the peak toward higher angles (causing a lattice shrinkage) and Mg toward lower angles (expanding the lattice). TEM investigations confirmed the precipitate-free microstructure of the milled powder. A similar process of precipitate dissolution during milling has been observed in mechanically alloyed Al7075 after 15 h of milling [47
], in a mechanically alloyed Al-Zn-Mg-Cu-Zr alloy after 40 h of milling [17
], in a spray-atomized 7075 alloy after 10 h of cryomilling [36
], and in a mechanically milled 7075 alloy after 20 h of milling [18
]. This extension of solubility can be attributed to enhanced diffusivity due to an increased dislocation density and volume fraction of atoms in the grain boundary area [5
The temperature of SPS (425 °C) is high enough to partially dissolve the phases Mg(Zn,Cu,Al)2
present in the cell boundaries of gas-atomized droplets [49
]. New precipitates are then formed during slow cooling from the SPS temperature. Due to a short thermal exposure during SPS, the solute atoms could not diffuse over a long distance and new precipitates are formed, especially at places where continuous layers of intermetallic phases were located in gas-atomized droplets, i.e., in cell boundaries. New precipitates are formed also at new boundaries formed during SPS between original droplets. The weight fraction of precipitates is higher in the SPS compact than in the gas-atomized droplets. This result is not surprising if the metastable state of the gas-atomized material is considered. On the other side, the temperature and especially the duration of SPS are not high enough for a remarkable precipitation of the Al3
Zr phase in the Al7075Zr_AC compact because of a very slow diffusion of Zr in the Al matrix [51
]. This is in agreement with experiments performed in the melt-spun AlZnMgCuZr alloy, which proved the formation of the L12
Zr phase during annealing for 1 h at a temperature as high as 460 °C [46
]. The presence of Al3
Zr particles was observed only by TEM investigations. Some grains contained homogeneously distributed Al3
Zr particles, while in other grains a fan-shaped arrangement was observed. This inhomogeneous distribution in individual grains is in correspondence with the literature.
Although the scheme of SPS was the same in both the atomized and milled powders, the resulting phase composition, and especially the distribution of precipitates, is different in both SPS compacts. The weight fraction of the main strengthening phase Mg(Zn,Cu,Al)2
is significantly higher in the Al7075Zr_MC compact. Similarly, only the Al7075Zr_MC compact contains the Al3
Zr phase in a volume fraction sufficient for detection by XRD. It can be supposed that the high density of dislocations and grain boundaries in the milled powder enhanced the diffusion processes in comparison with the atomized material so that the precipitation processes were accelerated. Intensive milling results not only in the dissolution of precipitates present in the gas-atomized droplets but also in a more homogeneous distribution of solutes. As a consequence, the precipitates formed during SPS are more homogeneously distributed. Further, the presence of Mg oxide, especially along the boundaries of the original powder particles, was also detected. It was shown in [9
] that even a small amount of Mg deoxidizes the amorphous surface layer of the Al oxide and forms new Mg-containing crystalline particles. The fine MgO particles can be also introduced into the powder particle interior during milling.
The microhardness of an atomized powder is determined predominantly by its phase composition. The 7xxx series Al alloys are known to exhibit natural aging at room temperature if prepared in a supersaturated state. The diffusion rate of the main alloying elements (Zn and Mg) is sufficiently high already at room temperature and redistribution of these solute atoms can occur. The atomized powder was stored after its preparation for a long time at room temperature, and it can be supposed that its phase composition is stable. The observed microhardness value of 97 ± 16 HV is deeply below the value achieved in the peak-aged ingot metallurgical Al7075 alloy (170 HV) [40
], and reflects the presence of coarse layers of intermetallic phases along the cell boundaries in the gas-atomized powder droplets. The same microhardness values have been found in the gas-atomized 7075 + 1 wt % Zr powder with a larger droplet size (98 ± 18 HV) [33
] and in the gas-atomized Al7075 powder (95 ± 18 HV) [21
]. This comparison shows negligible hardening effects of droplet size and Zr addition within the scatter of measured microhardness values. A similar low effect of Zr addition to the AlZnMgCu alloy was reported in [32
], where the microhardness of melt-spun ribbons was measured.
The milling of atomized powder particles introduced a large deformation energy into the material, and results in a grain refinement to the nanometer scale. Consequently, an extremely high microhardness value of 343 ± 25 HV was found. Although no optimization of the phase composition was performed, this value is about twofold of that observed in the peak-aged ingot metallurgical Al7075 alloy [40
] or in the melt-spun ribbons of the AlZnMgCu + 1 wt % Zr alloy after optimized thermal treatment (186 HV) [46
]. Figure 5
shows that the milled powder particles are contaminated by particles containing Ti and Ni. However, these particles are relatively coarse, and it is improbable that they could be responsible for the high microhardness. Additionally, as yet unpublished results obtained on the Al7075 alloy milled at the same conditions show a slightly lower microhardness despite a larger volume fraction of these contaminating particles present in this alloy. On the other side, the oxide particles introduced into the powder particles’ interiors during milling could strengthen the material [5
]. Surprisingly, the microhardness found in our milled material is significantly higher than that measured in the Al7075 alloy milled at cryotemperatures (220 ± 20 HV) using another milling device [34
]. This comparison suggests that the type of milling device has an important influence because of the different amounts of introduced deformation. The high efficiency of the mill used in our experiment can also explain the shorter times necessary for the complete dissolution of intermetallic phases in comparison with the literature [17
SPS influences significantly the strength of the materials. An increase in microhardness from 97 ± 16 HV up to 161 ± 6 HV was observed in the Al7075Zr_AC compact, while a decrease from 343 ± 25 HV to 195 ± 6 HV was found in the Al7075Zr_MC compact. These changes can be explained by an alteration in the microstructure and phase composition described above. In the Al7075Zr_AC compact, a partial dissolution occurring during sintering is followed by the formation of new precipitates during slow cooling from sintering temperature and during a long-term stay at room temperature between the sintering and microhardness measurement. Especially, the small Mg(Zn,Cu,Al)2
precipitates (Figure 16
a) are believed to produce the main hardening effect. In the Al7075Zr_MC compact, similar precipitation processes occur. On the other side, a recrystallization of the severely deformed microstructure of the milled powder occurred during SPS and resulted in a decrease in dislocation density and an increase in grain size (compare Figure 8
a and Figure 17
a); both of these processes reduce the material strength. Nevertheless, the microhardness is still close to the value of 200 HV, which is about 20% higher than that of the peak-aged ingot metallurgical counterpart. Especially, the sub-microcrystalline microstructure is believed to be responsible for this high microhardness.