- Email: [email protected]
Structure and mechanical properties of Al–Ca alloys processed by severe plastic deformation
Materials Science & Engineering A 767 (2019) 138410
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Structure and mechanical properties of Al–Ca alloys processed by severe plastic deformation
T
S.O. Rogacheva,∗, E.A. Naumovaa,b, E.S. Vasilevaa, M. Yu Magurinaa, R.V. Sundeeva,c, A.A. Veligzhanind The National University of Science and Technology “MISIS”, 4 Leninsky pr, Moscow, 119049, Russia MSTU “STANKIN”, 3a Vadkovskiy Pereulok, Moscow, 127055, Russia c Moscow Technological University “MIREA”, 78 Vernadskogo av, Moscow, 119454, Russia d National Research Centre “Kurchatov Institute”, 1 Akademika Kurchatova pl, Moscow, 123182, Russia a
b
A R T I C LE I N FO
A B S T R A C T
Keywords: Eutectic Al–Ca alloy High-pressure torsion Nanocrystalline materials Microstructure Hardening
The study investigated the effect of deformation by high-pressure torsion (HPT) on the hardening, structure transformations, and thermal stability of two eutectic alloys: binary – Al-8.0% Ca and complex-alloyed – Al-3.5% Ca-0.9% Mn-0.5% Fe-0.1% Zr-0.1% Sc. The HPT-deformation of the alloys (5 revolutions) led to the formation of a nanocrystalline structure with a high density of crystal defects. A predominant grain size was 20–40 nm in the binary alloy and 11–34 nm in the complex alloy. HPT resulted in the refinement of the Al4Ca particles for the binary alloy and the transformation of the Al4Ca particles into nanoclusters and segregation for the complex alloy. HPT increased the microhardness of the binary alloy to 1.80–2.05 GPa (~2 times), and of the complexalloyed alloy to 2.40–2.70 GPa (4.1–4.6 times). The hardening of the complex alloy is retained to higher heating temperatures compared to the binary alloy.
1. Introduction The wide use of aluminum based alloys as structural materials is associated with a unique combination of the most important performance properties (strength, ductility, corrosion resistance, electrical conductivity, etc.). At the same time, aluminum is characterized by low density and low cost (aluminum is the most abundant metal in the earth's crust) [1–3]. Recently, there have been studies of multiphase eutectic alloys based on the Al–Ca system with high technological properties both after casting and plastic deformation [4–8]. The structure of these alloys consists of an aluminum solid solution (Al) and eutectic [(Al)+Al4Ca], which has a very thin structure. Due to the low density of the Al4Ca intermetallic compound (2.33 g/cm3), these alloys have a low density. In work [4], high corrosion resistance of the aluminum-calcium alloys was mentioned, and these data confirmed recent studies [5–8]. The authors of [5] showed that the addition of iron in an amount of up to 1% does not adversely affect the structure and properties of the Al–Ca alloys, since iron is part of the Al10CaFe2 ternary eutectic intermetallic compound. In other aluminum alloys, iron is a harmful impurity due to the formation of the Al3Fe phase in the form of coarse needles [9]. Manganese, zirconium and scandium form a supersaturated solid ∗
solution in aluminum already during the casting process, which makes it possible to harden these alloys by annealing, without using the quenching operation [8]. Despite a number of the advantages described above, these Al–Ca alloys have a medium strength. Finding the possibility of improving their mechanical properties is of undoubted interest. The use of severe plastic deformations is a promising and rather well-developed method of the grain structure refinement and improving the complex properties of various metallic materials [10–14]. Numerous studies have shown that the use of severe plastic deformations can significantly increase the strength of materials such as pure aluminum and aluminum alloys [15–22]. The greatest refinement of the grain structure and the simultaneous extreme increase in strength is typical for the processing of aluminum alloys by high pressure torsion (HPT) [23–25]. In this work, the effect of the HPT-deformation on the structural transformations, hardening and thermal stability of two eutectic Al–Ca alloys (Al-8.0% Ca and Al-3.5% Ca-0.9% Mn-0.5% Fe-0.1% Zr-0.1% Sc) was studied.
Corresponding author. 4 Leninskiy ave, 119049, Moscow, Russia. E-mail address: [email protected] (S.O. Rogachev).
https://doi.org/10.1016/j.msea.2019.138410 Received 12 July 2019; Received in revised form 7 September 2019; Accepted 9 September 2019 Available online 10 September 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
Materials Science & Engineering A 767 (2019) 138410
S.O. Rogachev, et al.
2. Materials and research methods As the material for investigation, we used the cast eutectic aluminum alloys of the following chemical composition: binary Al-8.0% Ca and complex-alloyed Al-3.5% Ca-0.9% Mn-0.5% Fe-0.1% Zr-0.1% Sc [7]. The melting was carried out in an electric resistance furnace in a graphite-chased crucible based on high purity aluminium (99.99%). Calcium, manganese, iron, zirconium and scandium were introduced into the aluminium melt in the form of binary master alloys (Al–15%Ca, Al–10%Mn, Al–10%Fe, Al–15%Zr and Al-2%Sc, respectively). The casting was carried out in a graphite mould at a temperature of ~780 °C to obtain flat ingots with dimensions of 15 × 30 × 180 mm (the cooling rate during solidification was ~10 K/s). HPT was performed at room temperature, pressure of 6 GPa, and the number of revolutions N = 1 and 5 for the binary alloy and 1; 5 and 10 for the complex alloy using samples 8 mm in diameter and 0.7 mm in initial thickness. HPT were carried out under constrained conditions, i.e. the sample was placed in a 0.3 mm deep shaped hole located in the lower rotating anvil. After HPT, the thickness of the samples was about 0.4 mm. The uniformity of deformation after HPT was checked by measuring the Vickers microhardness (load 0.5 N, holding time 10 s) of the samples at two mutually perpendicular diameters with a step of 0.5 mm. Microhardness measurements were performed using a Micromet 5101 (Buehler) tester. Six measurements were made for each analyzed point (with a distance of 100 μm between adjacent indentations) with subsequent calculation of the average value. The microstructure of the samples was studied using a TESCAN VEGA scanning electron microscope and a Axio Observer D1m Carl Zeiss optical microscope. Electron-microscopic studies of thin foils were performed using a JEM-2100 transmission electron microscope (JEOL) with an accelerating voltage of 200 kV at magnifications up to × 50000. The foils were prepared by thinning the HPT-samples to a thickness of ~100 μm by mechanical grinding. Further, the disks with a diameter of 3 mm were cut out from the mid-radius of the thinned HPTsample and thinned by a jet electropolishing at a temperature of –20°С and a voltage of 20–30 V in electrolyte consisting of HNO3 in methanol. The chemical composition of the particles and aluminum matrix was determined using EDS micro-analysis. The transverse size of the structural elements was calculated using the ImageExpert software. At least 100 structural elements (grains, subgrains, particles) were measured for each state of the sample. Additionally, local chemical analysis of the samples was performed using a JSM-IT500 (JEOL) scanning electron microscopy with EDS. The study of the phase and structure composition of the samples was carried out by two methods:
Fig. 1. The distribution of microhardness values over the surface of the Al-8% Ca (a) and Al-3.5% Ca-0.9% Mn-0.5% Fe-0.1% Zr-0.1% Sc (b) alloy samples after HPT.
3. Results 3.1. Hardening of the alloys in the course of HPT The microhardness was 1.08 ± 0.03 GPa and 0.58 ± 0.03 GPa, respectively, for the binary and complex alloy for the original cast state. The HPT deformation led to a significant increase in the microhardness of the alloys and to the appearance of a non-uniform distribution of the microhardness values on the surface of the samples (Fig. 1). The minimum values of microhardness were observed in the centre of the samples, and the maximum – on their periphery. HPT (1 revolution) of the binary alloy led to an increase in microhardness values up to 1.80–2.05 GPa (~2 times) at the sample midradius, while at the sample centre the microhardness values did not change significantly (Fig. 1 a). An increase in the number of revolutions to N = 5 resulted in an almost uniform distribution of microhardness over the sample surface (1.70–2.05 GPa). HPT of the complex alloy through one revolution led to an increase in the microhardness values to 1.13–1.25 GPa at the sample mid-radius and to 0.83 ± 0.03 GPa at the sample centre (Fig. 1 b). An increase in the number of revolutions to N = 5 increased the microhardness values of the complex alloy to 2.40–2.70 GPa at the sample mid-radius and to 1.12 ± 0.1 GPa at the sample centre. Thus, HPT (5 revolutions)
- by X-ray diffractometry in a monochromatic CoKα radiation using a RIGAKU Ultima IV diffractometer. The analysis of the X-ray diffraction patterns and determination of the volume fraction of phases was performed using the PDXL program; - by X-ray diffraction method in synchrotron radiation in transmission mode with a spatial resolution of 350 μm. The measurement was performed in the area of the sample mid-radius, as well as at its periphery. To study of thermal stability, the samples after HPT were heated in an electric furnace in the temperature range from 100 to 400°С (with 50 °C step and 1 h exposure). Thermal stability was evaluated by the change in microhardness values measured at the sample mid-radius (six measurements for each measurement point with subsequent calculation of the average value).
2
Materials Science & Engineering A 767 (2019) 138410
S.O. Rogachev, et al.
Fig. 2. SEM-microstructure of the Al-8% Ca (a) and Al-3.5% Ca-0.9% Mn-0.5% Fe-0.1% Zr-0.1% Sc (b) alloys before HPT.
Al4Ca. According to the TEM data, it can be seen that an ultrafine-grained structure was formed in the complex alloy after HPT through one revolution (at the sample mid-radius) (Fig. 4 a). There are both large fragments (or subgrains) of irregular shape with a size of 300–450 nm, and smaller structural elements with a size of 50–300 nm. The high density of crystal defects is indicated by intragranular contrast. Individual point reflexes located along the rings on the electron diffraction pattern indicate the presence of high-angle boundaries. Some of these reflexes have azimuthal blurring, which indicates the presence of structural elements (fragments, subgrains) with low-angle boundaries. The presence of structural elements with low-angle and high-angle boundaries can be seen in the contrast on the TEM dark-field images. In the sample structure, the areas containing the ex-eutectic phase are visible, while the size of Al4Ca particles in these areas has decreased to 5–10 nm or less (Fig. 4 b). An increase in the number of revolutions to N = 5 led to the formation of a strongly deformed grain-subgrain nanocrystalline structure with a predominant size of structural elements of 11–34 nm for the complex alloy (at the sample mid-radius) (Fig. 4). Thus, the size of the structural elements in the complex alloy after HPT through 5 revolutions is an order of magnitude smaller than after HPT through 1 revolution. It can be seen that the diffraction rings on the diffraction pattern are formed by practically touching reflexes, which makes it possible to judge a more dispersed structure in comparison with the structure of the binary alloy. A less clear contrast on the TEM dark-field image in comparison with the binary alloy may indicate the presence of a structure with a larger share of structural elements with low-angle boundaries. EDS micro analysis of the microstructure of the alloy showed that the aluminum matrix is rich in calcium (from 0.8 to 3%). The electron diffraction pattern contains only reflexes of Al. The Al6(Mn, Fe) particles 0.5–2 μm (after 1 revolution) and 40–500 nm (after 5 revolutions) in size were also observed in the
resulted in an increase in the microhardness of the complex alloy by 4.1–4.6 times compared to the original cast state; at the same time, as the number of revolutions increased from 1 to 5, the nonuniform distribution of the microhardness values over the sample surface increased too. Increasing the number of revolutions to N = 10 resulted in a more uniform distribution of microhardness on the sample surface. It should be noted that in spite of the fact that HPT have a stronger effect on the increase in the microhardness of the complex alloy, compared to the binary one, the absolute values of microhardness after HPT for the complex alloy are only 1.2–1.5 times higher than for the binary alloy. 3.2. The microstructure of the alloys before and after HPT According to optical microscopy, the structure of both alloys in the original cast state consisted of large dendrites based on an aluminum (Al) with an average size of 6 ± 1 μm for the binary alloy and 16 ± 2 μm for the complex alloy and eutectic [(Al) + Al4Ca)] (Fig. 2). SEM analysis revealed a small amount of large particles enriched in iron and manganese in the structure of the complex alloy of the original state. Obviously, these are the primary precipitates of the Al6Mn phase, in which the manganese atoms are partially replaced by iron atoms, i.e. this phase can be described by the formula Al6(Mn, Fe). According to the TEM data, in the binary alloy, after HPT through 5 revolutions, a nanocrystalline structure formed with a large share of crystal defects and a predominant size of structural elements (grains, subgrains, fragments) of 20–40 nm (at the sample mid-radius) (Fig. 3). The presence of high-angle boundaries is indicated by numerous point reflections on the ring electron diffraction pattern and the contrast on the TEM dark field image. Many of these reflexes have azimuthal blurring, which indicates the presence of structural elements (fragments, subgrains) with low angle boundaries. The structure of the alloy also revealed a large number of small particles with a predominant size of 10–20 nm. According to EDS micro analysis, these particles are
Fig. 3. TEM-microstructure of the Al-8% Ca alloy after HPT through 5 revolutions: a – bright field; b – dark field in the Al-reflexes. 3
Materials Science & Engineering A 767 (2019) 138410
S.O. Rogachev, et al.
Fig. 4. TEM-microstructure of the Al-3.5% Ca-0.9% Mn-0.5% Fe-0.1% Zr-0.1% Sc alloy after HPT through 1 (a, b) and 5 (c, d) revolutions: a, b, c – the light field; d – dark field in the Al-reflexes.
temperature range of heating (150–300 °C), while for the complex alloy, a more complex dependence of the microhardness change under heating is observed. A sharp decrease in microhardness (1.7 times) is observed when heated to a temperature of 250 °C, after which the rate of decrease in microhardness values slows down when heated to a temperature of 300 °C, and with further heating in the temperature range of 350–400 °C, it is accelerated. The X-ray phase analysis revealed the Al3Zr and Al4Ca particles in the structure of the alloy heated to temperatures of 250–300 °C.
structure of the complex alloy processed by HPT. 3.3. The phase composition of the alloys before and after HPT According to the X-ray phase analysis data, the structure of the alloys in the original cast state consisted of (Al) matrix and eutectic [(Al) + Al4Ca] with the volume fraction of the Al4Ca phase ~ 29% and ~12%, respectively, for the binary and complex alloy. The intensity of the X-ray lines of the Al4Ca phase for the binary alloy after HPT does not change. At the same time, the intensity of the X-ray lines of the Al4Ca phase for the complex alloy is significantly weakened after just one revolution of HPT. HPT through 5 revolutions resulted in the complete disappearance of the Al4Ca lines. At the same time, the appearance of the lines of the Al6Mn type phase was observed. According to the data of quantitative X-ray phase analysis, the volume fraction of this phase was 2–3%. Obviously, this phase represents the Al6(Mn, Fe) particles, identified on the TEM images of the structure of the alloy after HPT. The results of the synchrotron phase analysis showed the presence of intense Al4Ca lines on the diffraction pattern for the binary alloy after HPT, both for the sample mid-radius and its periphery (Fig. 5 a, b). At the same time, very weak Al4Ca lines were observed for the complex alloy after HPT (Fig. 6 a). Moreover, their intensity tends to the background for the sample periphery (Fig. 6 b). For the binary alloy after HPT, no increase in the crystal lattice parameter of the solid aluminum solution was found, while for the complex alloy after HPT, the crystal lattice parameter of the solid aluminum solution increased by 0.13%.
4. Discussion of the results TEM analysis showed that the grain structure in both alloys is significantly refinement as a result of HPT; the Al4Ca eutectic particles are refinement in the binary alloy, and the Al4Ca eutectic particles are transformed into nanoclusters and segregation in the complex alloy. Nanoscale Al4Ca particles in the structure of the binary alloy after HPT are comparable to dispersed hardening particles of L12 type phases [7]. In both alloys a deformed grain-subgrain structure with a high density of crystal defects is formed as a result of HPT. In this case, in the complex alloy, the structure is more dispersed and is characterized by a large number of structural elements with low-angle boundaries and a higher density of crystal defects. This, apparently, is one of the reasons for the higher microhardness values of the complex alloy, as compared with the binary one. The formation of such structure in the complex alloy is due to the fact that the aluminum matrix of the alloy contains manganese, scandium and zirconium, and is characterized by a higher recrystallization temperature, therefore strain hardening as a result of the HPT deformation is more pronounced than for binary alloy. A significant decrease in the size of the Al6(Mn, Fe) particles in the structure of the complex alloy after HPT, compared with the original alloy, as well as the appearance of the corresponding peaks on the X-ray pattern, may indicate both fragmentation (refinement) of the initial
3.4. Thermal stability of the alloys after HPT Heating of the alloys processed by HPT to temperatures above 100 °C (for the binary alloy) and above 200 °C (for the complex alloy) resulted in a decrease in microhardness values (Fig. 7). For the binary alloy, the microhardness values decrease slowly over the entire 4
Materials Science & Engineering A 767 (2019) 138410
S.O. Rogachev, et al.
Fig. 5. Results of the synchrotron phase analysis for the Al-8% Ca alloy: (a) sample mid-radius; (b) sample periphery.
Fig. 6. Results of the synchrotron phase analysis for the Al-3.5% Ca-0.9% Mn0.5% Fe-0.1% Zr-0.1% Sc alloy: (a) sample mid-radius; (b) sample periphery.
large Al6(Mn, Fe) particles in the course of HPT, and precipitation of new smaller particles. Fragmentation or dissolution (precipitation) of the second phase particles was observed for many materials after severe plastic deformation [26–30]. In particular, fragmentation of the eutectic phase (Si) and the Al5FeSi phase and more homogeneous distribution in the aluminum matrix was observed in eutectic-containing Al-7% Si-0.3% Fe alloy in the course of the HPT process (room temperature, pressure 6 GPa, 5 revolutions) under constrained conditions [29]. On the contrary, dissolution of the Mg2Si, Al3Zr4 and AlZr3 phases was observed in the course of the ECAP process (room temperature, BC route, channel intersection angle 90°, number of passes N = 4) of the Al-0.16% Fe-0.51% Si-0.34% Mg-0.014% Mn-0.1% Zr and Al-0.16% Fe0.51% Si-0.34% Mg-0.014% Mn-0.1% Zr-0.1% Sc alloys [30]. Some authors attribute the possibility of phase transformations under process of severe plastic deformations with the value of the so-called “effective” temperature [31,32]. The inhomogeneous distribution of microhardness on the surface of the samples of aluminum alloys processed by HPT, which was observed in this work, was described earlier in a large number of works and is characteristic of many pure metals and alloys [33–36]. This is due, primarily, to the fact that, with a constant sample thickness, the degree of strain increases with increasing distance from the sample centre. This in turn leads to the formation of a different microstructure at the centre of the sample and at its periphery. An increase in the number of revolutions at HPT (i.e. the degree of strain) can lead to a decrease in the difference in microhardness values between the centre and the periphery of the sample due to the transfer of plastic flow of nearby material volumes. The authors of work [33] considered three models of the deformation behavior of a material under HPT: strain hardening, strain
softening and strain weakening. In general, it can be said that the implementation of a particular model depends on the ratio of the process of strain hardening, the process of formation of the developed grain structure and the process of dissolution of the second phases in the course of the HPT deformation. In our case, as the number of revolutions increased from 1 to 5, the difference in microhardness values between the centre of the sample and its periphery for the binary alloy decreased, and for the complex alloy, on the contrary, it increased. The detailed comparative studies of the complex alloy structure after 1 and 5 revolutions allowed to explain this behavior of the alloy. The hardening of the complex alloy in the course of HPT is influenced by the following factors: the strain hardening, the grain structure refinement and the transformation of the eutectic. After HPT through 1 revolution, the ultrafine grain-subgrain structure is formed and the eutectic is partially refinement, which leads to an increase in the microhardness values of ~2 times. A further significant increase in microhardness after 5 revolutions is mainly due to a significant decrease in the size of structural elements and the formation of a nanocrystalline deformed structure with a very high defect density (i.e., due strain hardening). In this case, the formation of such structure occurs primarily at the sample periphery, which leads to a significant increase in the microhardness values in these areas. In the centre of the sample, the eutectic areas are twisted around the centre of the sample (see Fig. 8 a). Accordingly, the structure in the sample centre is not being processing, which explains the low values of microhardness. With a further increase in the number of revolutions (to N = 10), the microhardness is equalized due to a more uniform processing of the structure in the sample volume. 5
Materials Science & Engineering A 767 (2019) 138410
S.O. Rogachev, et al.
The obtained results showed that, as a result of HPT, the complex alloy significantly reduces the size of Al4Ca particles, followed by the formation of nanoclusters and segregations. This, apparently, explains the weakening of the intensity of the X-ray lines from Al4Ca, and their disappearance with an increase in the number of the HPT revolutions. However, the generalized results of TEM, SEM and synchrotron analysis indirectly indicate the possibility of partial dissolution of the Al4Ca particles in the aluminum matrix of the complex alloy: the presence of only Al reflexes on the electron diffraction pattern, enrichment of the aluminum matrix with calcium at the level of the chemical composition of the alloy, an increase in the Al4Ca lattice parameter. The different nature of the change in microhardness during heating of the binary and complex alloys processed by HPT may be due to the following. Despite the lower (by ~ 100 °C) thermal stability of the binary alloy compared to the complex alloy, it has a slower rate of decrease in the microhardness values during heating. This is due to the presence of nano-sized Al4Ca particles, which slow down grain growth and couse dispersion hardening. The increased thermal stability of the complex alloy compared to the binary alloy may be associated with the doping of the aluminum matrix with manganese, zirconium, and scandium, which complicates the diffusion process and, consequently, grain growth during heating. The sharp drop in the microhardness of the complex alloy when heated to temperatures above 200 °C is probably associated with grain growth. The subsequent slowing down of the rate of the microhardness change of the complex alloy during heating in the temperature range 250–300 °C is associated with the release of Al3Zr and Al4Ca particles detected in the alloy structure by X-ray phase analysis. 5. Conclusion The results of the study of the HPT deformation effect on the structure transformations, hardening and thermal stability of two eutectic alloys: binary – Al-8.0% Ca and complex-alloyed – Al-3.5% Ca0.9% Mn-0.5% Fe-0.1% Zr-0.1% Sc showed:
Fig. 7. The effect of heating on the microhardness of the Al-8% Ca (a) and Al3.5% Ca-0.9% Mn-0.5% Fe-0.1% Zr-0.1% Sc (b) alloy samples processed by HPT.
(1) HPT (5 revolutions) resulted in the formation of a nanocrystalline deformed structure with a high density of crystal defects. A predominant grain size was 20–40 nm in the binary alloy and 11–34 nm in the complex alloy. The structure of the complex alloy is characterized by a large number of structural elements with lowangle boundaries and a higher density of crystal defects. (2) HPT led to refinement of the Al4Ca eutectic particles for the binary alloy (up to 10–20 nm) and transformed the Al4Ca eutectic particles into nanoclusters and segregation for the complex alloy. (3) HPT increased the microhardness of the binary alloy to 1.80–2.05 GPa (~2 times), and of the complex alloy to 2.40–2.70 GPa (4.1–4.6 times). The hardening of the complex alloy is maintained to higher heating temperatures, in comparison with
It should be noted that we were able to cut the foils for the TEM studies from the HPT-samples by the usual mechanical cutting without fracture of the HPT-samples, which suggests that the samples after HPT have both increased strength and some ductility (it is known that the foils for the TEM studies from brittle samples can be cut only by ultrasonic cutting, because mechanical cutting leads to the fracture of such samples). In further studies, it is advisable to increase the diameter of the samples for HPT to 20 mm. This will allow to make the specimens for tensile or fatigue tests and adequately evaluate the mechanical properties. These properties determine the possibility of practical use of the samples after HPT (e.g. for robotics).
Fig. 8. The surface of the centre (a, b) and mid-radius (c) of the samples of the Al-3.5% Ca-0.9% Mn-0.5% Fe-0.1% Zr-0.1% Sc alloy processed by HPT: (a) – 5 revolutions; (b, c) – 10 revolutions. 6
Materials Science & Engineering A 767 (2019) 138410
S.O. Rogachev, et al.
the binary alloy (up to 200 and 100°С, respectively).
Sci. Eng. A 519 (2009) 105–111 https://doi.org/10.1016/j.msea.2009.06.005. [17] M.H. Shaeri, M. Shaeri, M. Ebrahimi, M.T. Salehi, S.H. Seyyedein, Effect of ECAP temperature on microstructure and mechanical properties of Al–Zn–Mg–Cu alloy, Prog. Nat. Sci.: Mater. Int. 26 (2016) 182–191 https://doi.org/10.1016/j.pnsc. 2016.03.003. [18] G.J. Raab, R.Z. Valiev, T.C. Lowe, Y.T. Zhu, Continuous processing of ultrafine grained Al by ECAP-Conform, Mater. Sci. Eng. A 382 (2004) 30–34 https://doi.org/ 10.1016/j.msea.2004.04.021. [19] P. Leo, E. Cerri, P.P. De Marco, H.J. Roven, Properties and deformation behaviour of severe plastic deformed aluminium alloys, J. Mater. Process. Technol. 182 (2007) 207–214 https://doi.org/10.1016/j.jmatprotec.2006.07.038. [20] Z. Horita, T. Fujinami, M. Nemoto, T.G. Langdon, Equal-channel angular pressing of commercial aluminum alloys: grain refinement, thermal stability and tensile properties, Metall. Mater. Trans. A 31 (2000) 691–701 https://doi.org/10.1007/ s11661-000-0011-8. [21] S.A. Nikulin, S.V. Dobatkin, V.G. Khanzhin, S.O. Rogachev, S.A. Chakushin, Effect of submicrocrystalline structure and inclusions on the deformation and failure of aluminum alloys and titanium, Met. Sci. Heat Treat. 51 (2009) 208–217 https://doi. org/10.1007/s11041-009-9153-5. [22] G. Angella, P. Bassani, A. Tuissi, D. Ripamonti, M. Vedani, Microstructure evolution and aging kinetics of Al–Mg–Si and Al–Mg–Si–Sc alloys processed by ECAP, Mater. Sci. Forum 503–504 (2006) 493–498 https://doi.org/10.4028/www.scientific.net/ MSF.503-504.493. [23] A. Loucif, R.B. Figueiredo, T. Baudin, F. Brisset, T.G. Langdon, Microstructural evolution in an Al-6061 alloy processed by high-pressure torsion, Mater. Sci. Eng. A 527 (2010) 4864–4869 https://doi.org/10.1016/j.msea.2010.04.027. [24] K.S. Ghosh, N. Gao, M.J. Starink, Characterisation of high pressure torsion processed 7150 Al–Zn–Mg–Cu alloy, Mater. Sci. Eng. A 552 (2012) 164–171 https:// doi.org/10.1016/j.msea.2012.05.026. [25] H.-J. Lee, J.-K. Han, S. Janakiraman, B. Ahn, M. Kawasakia, T.G. Langdon, Significance of grain refinement on microstructure and mechanical properties of an Al-3% Mg alloy processed by high-pressure torsion, J. Alloy. Comp. 686 (2016) 998–1007 https://doi.org/10.1016/j.jallcom.2016.06.194. [26] S.O. Rogachev, S.A. Nikulin, V.M. Khatkevich, Structural and phase transformations in internally nitrided corrosion-resistant steel during severe plastic deformation and subsequent annealings, Steel Res. Int. 88 (2017) 1700070 https://doi.org/10.1002/ srin.201700070. [27] Yu Ivanisenko, W. Lojkowski, R.Z. Valiev, H.-J. Fecht, The mechanism of formation of nanostructure and dissolution of cementite in a pearlitic steel during high pressure torsion, Acta Mater. 51 (2003) 5555–5570 https://doi.org/10.1016/ S1359-6454(03)00419-1. [28] S.O. Rogachev, S.A. Nikulin, A.B. Rozhnov, M.V. Gorshenkov, Microstructure, phase composition, and thermal stability of two zirconium alloys subjected to highpressure torsion at different temperatures, Adv. Eng. Mater. 20 (2018) 1800151 https://doi.org/10.1002/adem.201800151. [29] C.M. Cepeda-Jiménez, J.M. García-Infanta, A.P. Zhilyaev, O.A. Ruano, F. Carreño, Influence of the thermal treatment on the deformation-induced precipitation of a hypoeutectic Al–7 wt% Si casting alloy deformed by high-pressure torsion, J. Alloy. Comp. 509 (2011) 636–643 https://doi.org/10.1016/j.jallcom.2010.09.122. [30] E. Cerri, P. Leo, Influence of severe plastic deformation on aging of Al–Mg–Si alloys, Mater. Sci. Eng. A 410–411 (2005) 226–229 https://doi.org/10.1016/j.msea.2005. 08.135. [31] B.B. Straumal, A.R. Kilmametov, Y. Ivanisenko, A.A. Mazilkin, O.A. Kogtenkova, L. Kurmanaeva, A. Korneva, P. Zięba, B. Baretzky, Phase transitions induced by severe plastic deformation: steady-state and equifinality, Int. J. Mater. Res. 106 (2015) 657–664 https://doi.org/10.3139/146.111215. [32] B. Straumal, A. Korneva, P. Zieba, Phase transitions in metallic alloys driven by the high pressure torsion, Arch. Civ. Mech. Eng. 14 (2014) 242–249 https://doi.org/10. 1016/j.acme.2013.07.002. [33] M. Kawasaki, R.B. Figueiredo, Y. Huang, T.G. Langdon, Interpretation of hardness evolution in metals processed by high-pressure torsion, J. Mater. Sci. 49 (2014) 6586–6596 https://doi.org/10.1007/s10853-014-8262-8. [34] A. Vorhauer, R. Pippan, On the homogeneity of deformation by high pressure torsion, Scr. Mater. 51 (2004) 921–925 https://doi.org/10.1016/j.scriptamat.2004.04. 025. [35] H. Jiang, Y.T. Zhu, D.P. Butt, I.V. Alexandrov, T.C. Lowe, Microstructural evolution, microhardness and thermal stability of HPT-processed Cu, Mater. Sci. Eng. A 290 (2000) 128–138 https://doi.org/10.1016/S0921-5093(00)00919-9. [36] Z. Yang, U. Welzel, Microstructure-microhardness relation of nanostructured Ni produced by high-pressure torsion, Mater. Lett. 59 (2005) 3406–3409 https://doi. org/10.1016/j.matlet.2005.05.077.
Data availability The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study. Acknowledgment The work was carried out with financial support from the Ministry of Science and Higher Education of the Russian Federation in the framework of Increase Competitiveness Program of NUST « MISiS» (№ К22019-008), implemented by a governmental decree dated 16th of March 2013, N 211. References [1] J.G. Kaufman, E.L. Rooy, Aluminum Alloy Castings: Properties, Processes and Applications, ASM International, Materials Park, USA, 2004. [2] G.S. Cole, A.M. Sherman, Lightweight materials for automotive applications, Mater. Char. 35 (1995) 3–9 https://doi.org/10.1016/1044-5803(95)00063-1. [3] V.S. Zolotorevskiy, N.A. Belov, Metal Science of Cast Aluminium Alloys, MISiS, Moscow, 2005 ([in russia]). [4] K. Swaminathan, K.A. Padmanabhan, Tensile flow and fracture behaviour of a superplastic Al–Ca–Zn alloy, J. Mater. Sci. 25 (1990) 4579–4586. [5] N.A. Belov, E.A. Naumova, V.D. Ilyukhin, V.V. Doroshenko, Structure and mechanical properties of Al – 6% Ca – 1% Fe alloy foundry goods, obtained by die casting, Tsvetnye Met. 3 (2017) 69–75 https://doi.org/10.17580/tsm.2017.03.11. [6] N.A. Belov, E.A. Naumova, T.K. Akopyan, Effect of calcium on structure, phase composition and hardening of Al–Zn–Mg alloys containing up to 12 wt.%Zn, Mater. Res. 18 (2015) 1384–1391 http://dx.doi.org/10.1590/1516-1439.036415. [7] N.A. Belov, K.A. Batyshev, V.V. Doroshenko, Microstructure and phase composition of the eutectic Al – Ca alloy, additionally alloyed with small additives of zirconium, scandium and manganese, Nonferrous Met. 2 (2017) 49–54 http://dx.doi.org/10. 17580/nfm.2017.02.09. [8] N.A. Belov, E.A. Naumova, A.N. Alabin, I.A. Matveeva, Effect of scandium on structure and hardening of Al–Ca eutectic alloys, J. Alloy. Comp. 646 (2015) 741–747 https://doi.org/10.1016/j.jallcom.2015.05.155. [9] D.G. Altenpohl, Aluminum: Technology, Applications, and Environment – a Profile of a Modern Metal, sixth ed., The Aluminum Association, Inc., Washington, D.C., 1998. [10] S.A. Nikulin, A.B. Rozhnov, S.O. Rogachev, V.M. Khatkevich, V.A. Turchenko, E.S. Khotulev, Investigation of structure, phase composition, and mechanical properties of Zr-2.5% Nb alloy after ECAP, Mater. Lett. 169 (2016) 223–226 https://doi.org/10.1016/j.matlet.2016.01.148. [11] R.Z. Valiev, Y. Estrin, Z. Horita, T.G. Langdon, M.J. Zehetbauer, Y.T. Zhu, Fundamentals of superior properties in bulk nanoSPD materials, Mater. Res. Lett. 4 (2016) 1–21 https://doi.org/10.1080/21663831.2015.1060543. [12] Y. Estrin, M.J. Zehetbauer, Niche applications of bulk nanostructured materials processed by severe plastic deformation, in: Michael J. Zehetbauer, Yuntian Theodore Zhu (Eds.), Bulk Nanostructured Materials, Wiley-VCH, Weinheim, Germany, 2009, pp. 635–648 https://doi.org/10.1002/9783527626892.ch28. [13] F.Y. Dong, P. Zhang, J.C. Pang, Y.B. Ren, K. Yang, Z.F. Zhang, Strength, damage and fracture behaviors of high-nitrogen austenitic stainless steel processed by highpressure torsion, Scr. Mater. 96 (2015) 5–8 https://doi.org/10.1016/j.scriptamat. 2014.09.016. [14] A.V. Polyakov, D.V. Gunderov, G.I. Raab, Evolution of microstructure and mechanical properties of titanium Grade 4 with the Increase of the ECAP-Conform passes, Mater. Sci. Forum 667–669 (2011) 1165–1170 https://doi.org/10.4028/ www.scientific.net/MSF.667-669.1165. [15] J. Estrin, M. Murashkin, R. Valiev, Ultrafine-grained aluminium alloys: processes, structural features and properties, in: R.N. Lumley (Ed.), Fundamentals of Aluminium Metallurgy, Woodhead Publishing, UK, 2011, pp. 468–503 https://doi. org/10.1533/9780857090256.2.468. [16] D. Orlov, Y. Beygelzimer, S. Synkov, V. Varyukhin, N. Tsuji, Z. Horita, Plastic flow, structure and mechanical properties in pure Al deformed by twist extrusion, Mater.
7
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Structure and mechanical properties of Al–Ca alloys processed by severe plastic deformation
T
S.O. Rogacheva,∗, E.A. Naumovaa,b, E.S. Vasilevaa, M. Yu Magurinaa, R.V. Sundeeva,c, A.A. Veligzhanind The National University of Science and Technology “MISIS”, 4 Leninsky pr, Moscow, 119049, Russia MSTU “STANKIN”, 3a Vadkovskiy Pereulok, Moscow, 127055, Russia c Moscow Technological University “MIREA”, 78 Vernadskogo av, Moscow, 119454, Russia d National Research Centre “Kurchatov Institute”, 1 Akademika Kurchatova pl, Moscow, 123182, Russia a
b
A R T I C LE I N FO
A B S T R A C T
Keywords: Eutectic Al–Ca alloy High-pressure torsion Nanocrystalline materials Microstructure Hardening
The study investigated the effect of deformation by high-pressure torsion (HPT) on the hardening, structure transformations, and thermal stability of two eutectic alloys: binary – Al-8.0% Ca and complex-alloyed – Al-3.5% Ca-0.9% Mn-0.5% Fe-0.1% Zr-0.1% Sc. The HPT-deformation of the alloys (5 revolutions) led to the formation of a nanocrystalline structure with a high density of crystal defects. A predominant grain size was 20–40 nm in the binary alloy and 11–34 nm in the complex alloy. HPT resulted in the refinement of the Al4Ca particles for the binary alloy and the transformation of the Al4Ca particles into nanoclusters and segregation for the complex alloy. HPT increased the microhardness of the binary alloy to 1.80–2.05 GPa (~2 times), and of the complexalloyed alloy to 2.40–2.70 GPa (4.1–4.6 times). The hardening of the complex alloy is retained to higher heating temperatures compared to the binary alloy.
1. Introduction The wide use of aluminum based alloys as structural materials is associated with a unique combination of the most important performance properties (strength, ductility, corrosion resistance, electrical conductivity, etc.). At the same time, aluminum is characterized by low density and low cost (aluminum is the most abundant metal in the earth's crust) [1–3]. Recently, there have been studies of multiphase eutectic alloys based on the Al–Ca system with high technological properties both after casting and plastic deformation [4–8]. The structure of these alloys consists of an aluminum solid solution (Al) and eutectic [(Al)+Al4Ca], which has a very thin structure. Due to the low density of the Al4Ca intermetallic compound (2.33 g/cm3), these alloys have a low density. In work [4], high corrosion resistance of the aluminum-calcium alloys was mentioned, and these data confirmed recent studies [5–8]. The authors of [5] showed that the addition of iron in an amount of up to 1% does not adversely affect the structure and properties of the Al–Ca alloys, since iron is part of the Al10CaFe2 ternary eutectic intermetallic compound. In other aluminum alloys, iron is a harmful impurity due to the formation of the Al3Fe phase in the form of coarse needles [9]. Manganese, zirconium and scandium form a supersaturated solid ∗
solution in aluminum already during the casting process, which makes it possible to harden these alloys by annealing, without using the quenching operation [8]. Despite a number of the advantages described above, these Al–Ca alloys have a medium strength. Finding the possibility of improving their mechanical properties is of undoubted interest. The use of severe plastic deformations is a promising and rather well-developed method of the grain structure refinement and improving the complex properties of various metallic materials [10–14]. Numerous studies have shown that the use of severe plastic deformations can significantly increase the strength of materials such as pure aluminum and aluminum alloys [15–22]. The greatest refinement of the grain structure and the simultaneous extreme increase in strength is typical for the processing of aluminum alloys by high pressure torsion (HPT) [23–25]. In this work, the effect of the HPT-deformation on the structural transformations, hardening and thermal stability of two eutectic Al–Ca alloys (Al-8.0% Ca and Al-3.5% Ca-0.9% Mn-0.5% Fe-0.1% Zr-0.1% Sc) was studied.
Corresponding author. 4 Leninskiy ave, 119049, Moscow, Russia. E-mail address: [email protected] (S.O. Rogachev).
https://doi.org/10.1016/j.msea.2019.138410 Received 12 July 2019; Received in revised form 7 September 2019; Accepted 9 September 2019 Available online 10 September 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
Materials Science & Engineering A 767 (2019) 138410
S.O. Rogachev, et al.
2. Materials and research methods As the material for investigation, we used the cast eutectic aluminum alloys of the following chemical composition: binary Al-8.0% Ca and complex-alloyed Al-3.5% Ca-0.9% Mn-0.5% Fe-0.1% Zr-0.1% Sc [7]. The melting was carried out in an electric resistance furnace in a graphite-chased crucible based on high purity aluminium (99.99%). Calcium, manganese, iron, zirconium and scandium were introduced into the aluminium melt in the form of binary master alloys (Al–15%Ca, Al–10%Mn, Al–10%Fe, Al–15%Zr and Al-2%Sc, respectively). The casting was carried out in a graphite mould at a temperature of ~780 °C to obtain flat ingots with dimensions of 15 × 30 × 180 mm (the cooling rate during solidification was ~10 K/s). HPT was performed at room temperature, pressure of 6 GPa, and the number of revolutions N = 1 and 5 for the binary alloy and 1; 5 and 10 for the complex alloy using samples 8 mm in diameter and 0.7 mm in initial thickness. HPT were carried out under constrained conditions, i.e. the sample was placed in a 0.3 mm deep shaped hole located in the lower rotating anvil. After HPT, the thickness of the samples was about 0.4 mm. The uniformity of deformation after HPT was checked by measuring the Vickers microhardness (load 0.5 N, holding time 10 s) of the samples at two mutually perpendicular diameters with a step of 0.5 mm. Microhardness measurements were performed using a Micromet 5101 (Buehler) tester. Six measurements were made for each analyzed point (with a distance of 100 μm between adjacent indentations) with subsequent calculation of the average value. The microstructure of the samples was studied using a TESCAN VEGA scanning electron microscope and a Axio Observer D1m Carl Zeiss optical microscope. Electron-microscopic studies of thin foils were performed using a JEM-2100 transmission electron microscope (JEOL) with an accelerating voltage of 200 kV at magnifications up to × 50000. The foils were prepared by thinning the HPT-samples to a thickness of ~100 μm by mechanical grinding. Further, the disks with a diameter of 3 mm were cut out from the mid-radius of the thinned HPTsample and thinned by a jet electropolishing at a temperature of –20°С and a voltage of 20–30 V in electrolyte consisting of HNO3 in methanol. The chemical composition of the particles and aluminum matrix was determined using EDS micro-analysis. The transverse size of the structural elements was calculated using the ImageExpert software. At least 100 structural elements (grains, subgrains, particles) were measured for each state of the sample. Additionally, local chemical analysis of the samples was performed using a JSM-IT500 (JEOL) scanning electron microscopy with EDS. The study of the phase and structure composition of the samples was carried out by two methods:
Fig. 1. The distribution of microhardness values over the surface of the Al-8% Ca (a) and Al-3.5% Ca-0.9% Mn-0.5% Fe-0.1% Zr-0.1% Sc (b) alloy samples after HPT.
3. Results 3.1. Hardening of the alloys in the course of HPT The microhardness was 1.08 ± 0.03 GPa and 0.58 ± 0.03 GPa, respectively, for the binary and complex alloy for the original cast state. The HPT deformation led to a significant increase in the microhardness of the alloys and to the appearance of a non-uniform distribution of the microhardness values on the surface of the samples (Fig. 1). The minimum values of microhardness were observed in the centre of the samples, and the maximum – on their periphery. HPT (1 revolution) of the binary alloy led to an increase in microhardness values up to 1.80–2.05 GPa (~2 times) at the sample midradius, while at the sample centre the microhardness values did not change significantly (Fig. 1 a). An increase in the number of revolutions to N = 5 resulted in an almost uniform distribution of microhardness over the sample surface (1.70–2.05 GPa). HPT of the complex alloy through one revolution led to an increase in the microhardness values to 1.13–1.25 GPa at the sample mid-radius and to 0.83 ± 0.03 GPa at the sample centre (Fig. 1 b). An increase in the number of revolutions to N = 5 increased the microhardness values of the complex alloy to 2.40–2.70 GPa at the sample mid-radius and to 1.12 ± 0.1 GPa at the sample centre. Thus, HPT (5 revolutions)
- by X-ray diffractometry in a monochromatic CoKα radiation using a RIGAKU Ultima IV diffractometer. The analysis of the X-ray diffraction patterns and determination of the volume fraction of phases was performed using the PDXL program; - by X-ray diffraction method in synchrotron radiation in transmission mode with a spatial resolution of 350 μm. The measurement was performed in the area of the sample mid-radius, as well as at its periphery. To study of thermal stability, the samples after HPT were heated in an electric furnace in the temperature range from 100 to 400°С (with 50 °C step and 1 h exposure). Thermal stability was evaluated by the change in microhardness values measured at the sample mid-radius (six measurements for each measurement point with subsequent calculation of the average value).
2
Materials Science & Engineering A 767 (2019) 138410
S.O. Rogachev, et al.
Fig. 2. SEM-microstructure of the Al-8% Ca (a) and Al-3.5% Ca-0.9% Mn-0.5% Fe-0.1% Zr-0.1% Sc (b) alloys before HPT.
Al4Ca. According to the TEM data, it can be seen that an ultrafine-grained structure was formed in the complex alloy after HPT through one revolution (at the sample mid-radius) (Fig. 4 a). There are both large fragments (or subgrains) of irregular shape with a size of 300–450 nm, and smaller structural elements with a size of 50–300 nm. The high density of crystal defects is indicated by intragranular contrast. Individual point reflexes located along the rings on the electron diffraction pattern indicate the presence of high-angle boundaries. Some of these reflexes have azimuthal blurring, which indicates the presence of structural elements (fragments, subgrains) with low-angle boundaries. The presence of structural elements with low-angle and high-angle boundaries can be seen in the contrast on the TEM dark-field images. In the sample structure, the areas containing the ex-eutectic phase are visible, while the size of Al4Ca particles in these areas has decreased to 5–10 nm or less (Fig. 4 b). An increase in the number of revolutions to N = 5 led to the formation of a strongly deformed grain-subgrain nanocrystalline structure with a predominant size of structural elements of 11–34 nm for the complex alloy (at the sample mid-radius) (Fig. 4). Thus, the size of the structural elements in the complex alloy after HPT through 5 revolutions is an order of magnitude smaller than after HPT through 1 revolution. It can be seen that the diffraction rings on the diffraction pattern are formed by practically touching reflexes, which makes it possible to judge a more dispersed structure in comparison with the structure of the binary alloy. A less clear contrast on the TEM dark-field image in comparison with the binary alloy may indicate the presence of a structure with a larger share of structural elements with low-angle boundaries. EDS micro analysis of the microstructure of the alloy showed that the aluminum matrix is rich in calcium (from 0.8 to 3%). The electron diffraction pattern contains only reflexes of Al. The Al6(Mn, Fe) particles 0.5–2 μm (after 1 revolution) and 40–500 nm (after 5 revolutions) in size were also observed in the
resulted in an increase in the microhardness of the complex alloy by 4.1–4.6 times compared to the original cast state; at the same time, as the number of revolutions increased from 1 to 5, the nonuniform distribution of the microhardness values over the sample surface increased too. Increasing the number of revolutions to N = 10 resulted in a more uniform distribution of microhardness on the sample surface. It should be noted that in spite of the fact that HPT have a stronger effect on the increase in the microhardness of the complex alloy, compared to the binary one, the absolute values of microhardness after HPT for the complex alloy are only 1.2–1.5 times higher than for the binary alloy. 3.2. The microstructure of the alloys before and after HPT According to optical microscopy, the structure of both alloys in the original cast state consisted of large dendrites based on an aluminum (Al) with an average size of 6 ± 1 μm for the binary alloy and 16 ± 2 μm for the complex alloy and eutectic [(Al) + Al4Ca)] (Fig. 2). SEM analysis revealed a small amount of large particles enriched in iron and manganese in the structure of the complex alloy of the original state. Obviously, these are the primary precipitates of the Al6Mn phase, in which the manganese atoms are partially replaced by iron atoms, i.e. this phase can be described by the formula Al6(Mn, Fe). According to the TEM data, in the binary alloy, after HPT through 5 revolutions, a nanocrystalline structure formed with a large share of crystal defects and a predominant size of structural elements (grains, subgrains, fragments) of 20–40 nm (at the sample mid-radius) (Fig. 3). The presence of high-angle boundaries is indicated by numerous point reflections on the ring electron diffraction pattern and the contrast on the TEM dark field image. Many of these reflexes have azimuthal blurring, which indicates the presence of structural elements (fragments, subgrains) with low angle boundaries. The structure of the alloy also revealed a large number of small particles with a predominant size of 10–20 nm. According to EDS micro analysis, these particles are
Fig. 3. TEM-microstructure of the Al-8% Ca alloy after HPT through 5 revolutions: a – bright field; b – dark field in the Al-reflexes. 3
Materials Science & Engineering A 767 (2019) 138410
S.O. Rogachev, et al.
Fig. 4. TEM-microstructure of the Al-3.5% Ca-0.9% Mn-0.5% Fe-0.1% Zr-0.1% Sc alloy after HPT through 1 (a, b) and 5 (c, d) revolutions: a, b, c – the light field; d – dark field in the Al-reflexes.
temperature range of heating (150–300 °C), while for the complex alloy, a more complex dependence of the microhardness change under heating is observed. A sharp decrease in microhardness (1.7 times) is observed when heated to a temperature of 250 °C, after which the rate of decrease in microhardness values slows down when heated to a temperature of 300 °C, and with further heating in the temperature range of 350–400 °C, it is accelerated. The X-ray phase analysis revealed the Al3Zr and Al4Ca particles in the structure of the alloy heated to temperatures of 250–300 °C.
structure of the complex alloy processed by HPT. 3.3. The phase composition of the alloys before and after HPT According to the X-ray phase analysis data, the structure of the alloys in the original cast state consisted of (Al) matrix and eutectic [(Al) + Al4Ca] with the volume fraction of the Al4Ca phase ~ 29% and ~12%, respectively, for the binary and complex alloy. The intensity of the X-ray lines of the Al4Ca phase for the binary alloy after HPT does not change. At the same time, the intensity of the X-ray lines of the Al4Ca phase for the complex alloy is significantly weakened after just one revolution of HPT. HPT through 5 revolutions resulted in the complete disappearance of the Al4Ca lines. At the same time, the appearance of the lines of the Al6Mn type phase was observed. According to the data of quantitative X-ray phase analysis, the volume fraction of this phase was 2–3%. Obviously, this phase represents the Al6(Mn, Fe) particles, identified on the TEM images of the structure of the alloy after HPT. The results of the synchrotron phase analysis showed the presence of intense Al4Ca lines on the diffraction pattern for the binary alloy after HPT, both for the sample mid-radius and its periphery (Fig. 5 a, b). At the same time, very weak Al4Ca lines were observed for the complex alloy after HPT (Fig. 6 a). Moreover, their intensity tends to the background for the sample periphery (Fig. 6 b). For the binary alloy after HPT, no increase in the crystal lattice parameter of the solid aluminum solution was found, while for the complex alloy after HPT, the crystal lattice parameter of the solid aluminum solution increased by 0.13%.
4. Discussion of the results TEM analysis showed that the grain structure in both alloys is significantly refinement as a result of HPT; the Al4Ca eutectic particles are refinement in the binary alloy, and the Al4Ca eutectic particles are transformed into nanoclusters and segregation in the complex alloy. Nanoscale Al4Ca particles in the structure of the binary alloy after HPT are comparable to dispersed hardening particles of L12 type phases [7]. In both alloys a deformed grain-subgrain structure with a high density of crystal defects is formed as a result of HPT. In this case, in the complex alloy, the structure is more dispersed and is characterized by a large number of structural elements with low-angle boundaries and a higher density of crystal defects. This, apparently, is one of the reasons for the higher microhardness values of the complex alloy, as compared with the binary one. The formation of such structure in the complex alloy is due to the fact that the aluminum matrix of the alloy contains manganese, scandium and zirconium, and is characterized by a higher recrystallization temperature, therefore strain hardening as a result of the HPT deformation is more pronounced than for binary alloy. A significant decrease in the size of the Al6(Mn, Fe) particles in the structure of the complex alloy after HPT, compared with the original alloy, as well as the appearance of the corresponding peaks on the X-ray pattern, may indicate both fragmentation (refinement) of the initial
3.4. Thermal stability of the alloys after HPT Heating of the alloys processed by HPT to temperatures above 100 °C (for the binary alloy) and above 200 °C (for the complex alloy) resulted in a decrease in microhardness values (Fig. 7). For the binary alloy, the microhardness values decrease slowly over the entire 4
Materials Science & Engineering A 767 (2019) 138410
S.O. Rogachev, et al.
Fig. 5. Results of the synchrotron phase analysis for the Al-8% Ca alloy: (a) sample mid-radius; (b) sample periphery.
Fig. 6. Results of the synchrotron phase analysis for the Al-3.5% Ca-0.9% Mn0.5% Fe-0.1% Zr-0.1% Sc alloy: (a) sample mid-radius; (b) sample periphery.
large Al6(Mn, Fe) particles in the course of HPT, and precipitation of new smaller particles. Fragmentation or dissolution (precipitation) of the second phase particles was observed for many materials after severe plastic deformation [26–30]. In particular, fragmentation of the eutectic phase (Si) and the Al5FeSi phase and more homogeneous distribution in the aluminum matrix was observed in eutectic-containing Al-7% Si-0.3% Fe alloy in the course of the HPT process (room temperature, pressure 6 GPa, 5 revolutions) under constrained conditions [29]. On the contrary, dissolution of the Mg2Si, Al3Zr4 and AlZr3 phases was observed in the course of the ECAP process (room temperature, BC route, channel intersection angle 90°, number of passes N = 4) of the Al-0.16% Fe-0.51% Si-0.34% Mg-0.014% Mn-0.1% Zr and Al-0.16% Fe0.51% Si-0.34% Mg-0.014% Mn-0.1% Zr-0.1% Sc alloys [30]. Some authors attribute the possibility of phase transformations under process of severe plastic deformations with the value of the so-called “effective” temperature [31,32]. The inhomogeneous distribution of microhardness on the surface of the samples of aluminum alloys processed by HPT, which was observed in this work, was described earlier in a large number of works and is characteristic of many pure metals and alloys [33–36]. This is due, primarily, to the fact that, with a constant sample thickness, the degree of strain increases with increasing distance from the sample centre. This in turn leads to the formation of a different microstructure at the centre of the sample and at its periphery. An increase in the number of revolutions at HPT (i.e. the degree of strain) can lead to a decrease in the difference in microhardness values between the centre and the periphery of the sample due to the transfer of plastic flow of nearby material volumes. The authors of work [33] considered three models of the deformation behavior of a material under HPT: strain hardening, strain
softening and strain weakening. In general, it can be said that the implementation of a particular model depends on the ratio of the process of strain hardening, the process of formation of the developed grain structure and the process of dissolution of the second phases in the course of the HPT deformation. In our case, as the number of revolutions increased from 1 to 5, the difference in microhardness values between the centre of the sample and its periphery for the binary alloy decreased, and for the complex alloy, on the contrary, it increased. The detailed comparative studies of the complex alloy structure after 1 and 5 revolutions allowed to explain this behavior of the alloy. The hardening of the complex alloy in the course of HPT is influenced by the following factors: the strain hardening, the grain structure refinement and the transformation of the eutectic. After HPT through 1 revolution, the ultrafine grain-subgrain structure is formed and the eutectic is partially refinement, which leads to an increase in the microhardness values of ~2 times. A further significant increase in microhardness after 5 revolutions is mainly due to a significant decrease in the size of structural elements and the formation of a nanocrystalline deformed structure with a very high defect density (i.e., due strain hardening). In this case, the formation of such structure occurs primarily at the sample periphery, which leads to a significant increase in the microhardness values in these areas. In the centre of the sample, the eutectic areas are twisted around the centre of the sample (see Fig. 8 a). Accordingly, the structure in the sample centre is not being processing, which explains the low values of microhardness. With a further increase in the number of revolutions (to N = 10), the microhardness is equalized due to a more uniform processing of the structure in the sample volume. 5
Materials Science & Engineering A 767 (2019) 138410
S.O. Rogachev, et al.
The obtained results showed that, as a result of HPT, the complex alloy significantly reduces the size of Al4Ca particles, followed by the formation of nanoclusters and segregations. This, apparently, explains the weakening of the intensity of the X-ray lines from Al4Ca, and their disappearance with an increase in the number of the HPT revolutions. However, the generalized results of TEM, SEM and synchrotron analysis indirectly indicate the possibility of partial dissolution of the Al4Ca particles in the aluminum matrix of the complex alloy: the presence of only Al reflexes on the electron diffraction pattern, enrichment of the aluminum matrix with calcium at the level of the chemical composition of the alloy, an increase in the Al4Ca lattice parameter. The different nature of the change in microhardness during heating of the binary and complex alloys processed by HPT may be due to the following. Despite the lower (by ~ 100 °C) thermal stability of the binary alloy compared to the complex alloy, it has a slower rate of decrease in the microhardness values during heating. This is due to the presence of nano-sized Al4Ca particles, which slow down grain growth and couse dispersion hardening. The increased thermal stability of the complex alloy compared to the binary alloy may be associated with the doping of the aluminum matrix with manganese, zirconium, and scandium, which complicates the diffusion process and, consequently, grain growth during heating. The sharp drop in the microhardness of the complex alloy when heated to temperatures above 200 °C is probably associated with grain growth. The subsequent slowing down of the rate of the microhardness change of the complex alloy during heating in the temperature range 250–300 °C is associated with the release of Al3Zr and Al4Ca particles detected in the alloy structure by X-ray phase analysis. 5. Conclusion The results of the study of the HPT deformation effect on the structure transformations, hardening and thermal stability of two eutectic alloys: binary – Al-8.0% Ca and complex-alloyed – Al-3.5% Ca0.9% Mn-0.5% Fe-0.1% Zr-0.1% Sc showed:
Fig. 7. The effect of heating on the microhardness of the Al-8% Ca (a) and Al3.5% Ca-0.9% Mn-0.5% Fe-0.1% Zr-0.1% Sc (b) alloy samples processed by HPT.
(1) HPT (5 revolutions) resulted in the formation of a nanocrystalline deformed structure with a high density of crystal defects. A predominant grain size was 20–40 nm in the binary alloy and 11–34 nm in the complex alloy. The structure of the complex alloy is characterized by a large number of structural elements with lowangle boundaries and a higher density of crystal defects. (2) HPT led to refinement of the Al4Ca eutectic particles for the binary alloy (up to 10–20 nm) and transformed the Al4Ca eutectic particles into nanoclusters and segregation for the complex alloy. (3) HPT increased the microhardness of the binary alloy to 1.80–2.05 GPa (~2 times), and of the complex alloy to 2.40–2.70 GPa (4.1–4.6 times). The hardening of the complex alloy is maintained to higher heating temperatures, in comparison with
It should be noted that we were able to cut the foils for the TEM studies from the HPT-samples by the usual mechanical cutting without fracture of the HPT-samples, which suggests that the samples after HPT have both increased strength and some ductility (it is known that the foils for the TEM studies from brittle samples can be cut only by ultrasonic cutting, because mechanical cutting leads to the fracture of such samples). In further studies, it is advisable to increase the diameter of the samples for HPT to 20 mm. This will allow to make the specimens for tensile or fatigue tests and adequately evaluate the mechanical properties. These properties determine the possibility of practical use of the samples after HPT (e.g. for robotics).
Fig. 8. The surface of the centre (a, b) and mid-radius (c) of the samples of the Al-3.5% Ca-0.9% Mn-0.5% Fe-0.1% Zr-0.1% Sc alloy processed by HPT: (a) – 5 revolutions; (b, c) – 10 revolutions. 6
Materials Science & Engineering A 767 (2019) 138410
S.O. Rogachev, et al.
the binary alloy (up to 200 and 100°С, respectively).
Sci. Eng. A 519 (2009) 105–111 https://doi.org/10.1016/j.msea.2009.06.005. [17] M.H. Shaeri, M. Shaeri, M. Ebrahimi, M.T. Salehi, S.H. Seyyedein, Effect of ECAP temperature on microstructure and mechanical properties of Al–Zn–Mg–Cu alloy, Prog. Nat. Sci.: Mater. Int. 26 (2016) 182–191 https://doi.org/10.1016/j.pnsc. 2016.03.003. [18] G.J. Raab, R.Z. Valiev, T.C. Lowe, Y.T. Zhu, Continuous processing of ultrafine grained Al by ECAP-Conform, Mater. Sci. Eng. A 382 (2004) 30–34 https://doi.org/ 10.1016/j.msea.2004.04.021. [19] P. Leo, E. Cerri, P.P. De Marco, H.J. Roven, Properties and deformation behaviour of severe plastic deformed aluminium alloys, J. Mater. Process. Technol. 182 (2007) 207–214 https://doi.org/10.1016/j.jmatprotec.2006.07.038. [20] Z. Horita, T. Fujinami, M. Nemoto, T.G. Langdon, Equal-channel angular pressing of commercial aluminum alloys: grain refinement, thermal stability and tensile properties, Metall. Mater. Trans. A 31 (2000) 691–701 https://doi.org/10.1007/ s11661-000-0011-8. [21] S.A. Nikulin, S.V. Dobatkin, V.G. Khanzhin, S.O. Rogachev, S.A. Chakushin, Effect of submicrocrystalline structure and inclusions on the deformation and failure of aluminum alloys and titanium, Met. Sci. Heat Treat. 51 (2009) 208–217 https://doi. org/10.1007/s11041-009-9153-5. [22] G. Angella, P. Bassani, A. Tuissi, D. Ripamonti, M. Vedani, Microstructure evolution and aging kinetics of Al–Mg–Si and Al–Mg–Si–Sc alloys processed by ECAP, Mater. Sci. Forum 503–504 (2006) 493–498 https://doi.org/10.4028/www.scientific.net/ MSF.503-504.493. [23] A. Loucif, R.B. Figueiredo, T. Baudin, F. Brisset, T.G. Langdon, Microstructural evolution in an Al-6061 alloy processed by high-pressure torsion, Mater. Sci. Eng. A 527 (2010) 4864–4869 https://doi.org/10.1016/j.msea.2010.04.027. [24] K.S. Ghosh, N. Gao, M.J. Starink, Characterisation of high pressure torsion processed 7150 Al–Zn–Mg–Cu alloy, Mater. Sci. Eng. A 552 (2012) 164–171 https:// doi.org/10.1016/j.msea.2012.05.026. [25] H.-J. Lee, J.-K. Han, S. Janakiraman, B. Ahn, M. Kawasakia, T.G. Langdon, Significance of grain refinement on microstructure and mechanical properties of an Al-3% Mg alloy processed by high-pressure torsion, J. Alloy. Comp. 686 (2016) 998–1007 https://doi.org/10.1016/j.jallcom.2016.06.194. [26] S.O. Rogachev, S.A. Nikulin, V.M. Khatkevich, Structural and phase transformations in internally nitrided corrosion-resistant steel during severe plastic deformation and subsequent annealings, Steel Res. Int. 88 (2017) 1700070 https://doi.org/10.1002/ srin.201700070. [27] Yu Ivanisenko, W. Lojkowski, R.Z. Valiev, H.-J. Fecht, The mechanism of formation of nanostructure and dissolution of cementite in a pearlitic steel during high pressure torsion, Acta Mater. 51 (2003) 5555–5570 https://doi.org/10.1016/ S1359-6454(03)00419-1. [28] S.O. Rogachev, S.A. Nikulin, A.B. Rozhnov, M.V. Gorshenkov, Microstructure, phase composition, and thermal stability of two zirconium alloys subjected to highpressure torsion at different temperatures, Adv. Eng. Mater. 20 (2018) 1800151 https://doi.org/10.1002/adem.201800151. [29] C.M. Cepeda-Jiménez, J.M. García-Infanta, A.P. Zhilyaev, O.A. Ruano, F. Carreño, Influence of the thermal treatment on the deformation-induced precipitation of a hypoeutectic Al–7 wt% Si casting alloy deformed by high-pressure torsion, J. Alloy. Comp. 509 (2011) 636–643 https://doi.org/10.1016/j.jallcom.2010.09.122. [30] E. Cerri, P. Leo, Influence of severe plastic deformation on aging of Al–Mg–Si alloys, Mater. Sci. Eng. A 410–411 (2005) 226–229 https://doi.org/10.1016/j.msea.2005. 08.135. [31] B.B. Straumal, A.R. Kilmametov, Y. Ivanisenko, A.A. Mazilkin, O.A. Kogtenkova, L. Kurmanaeva, A. Korneva, P. Zięba, B. Baretzky, Phase transitions induced by severe plastic deformation: steady-state and equifinality, Int. J. Mater. Res. 106 (2015) 657–664 https://doi.org/10.3139/146.111215. [32] B. Straumal, A. Korneva, P. Zieba, Phase transitions in metallic alloys driven by the high pressure torsion, Arch. Civ. Mech. Eng. 14 (2014) 242–249 https://doi.org/10. 1016/j.acme.2013.07.002. [33] M. Kawasaki, R.B. Figueiredo, Y. Huang, T.G. Langdon, Interpretation of hardness evolution in metals processed by high-pressure torsion, J. Mater. Sci. 49 (2014) 6586–6596 https://doi.org/10.1007/s10853-014-8262-8. [34] A. Vorhauer, R. Pippan, On the homogeneity of deformation by high pressure torsion, Scr. Mater. 51 (2004) 921–925 https://doi.org/10.1016/j.scriptamat.2004.04. 025. [35] H. Jiang, Y.T. Zhu, D.P. Butt, I.V. Alexandrov, T.C. Lowe, Microstructural evolution, microhardness and thermal stability of HPT-processed Cu, Mater. Sci. Eng. A 290 (2000) 128–138 https://doi.org/10.1016/S0921-5093(00)00919-9. [36] Z. Yang, U. Welzel, Microstructure-microhardness relation of nanostructured Ni produced by high-pressure torsion, Mater. Lett. 59 (2005) 3406–3409 https://doi. org/10.1016/j.matlet.2005.05.077.
Data availability The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study. Acknowledgment The work was carried out with financial support from the Ministry of Science and Higher Education of the Russian Federation in the framework of Increase Competitiveness Program of NUST « MISiS» (№ К22019-008), implemented by a governmental decree dated 16th of March 2013, N 211. References [1] J.G. Kaufman, E.L. Rooy, Aluminum Alloy Castings: Properties, Processes and Applications, ASM International, Materials Park, USA, 2004. [2] G.S. Cole, A.M. Sherman, Lightweight materials for automotive applications, Mater. Char. 35 (1995) 3–9 https://doi.org/10.1016/1044-5803(95)00063-1. [3] V.S. Zolotorevskiy, N.A. Belov, Metal Science of Cast Aluminium Alloys, MISiS, Moscow, 2005 ([in russia]). [4] K. Swaminathan, K.A. Padmanabhan, Tensile flow and fracture behaviour of a superplastic Al–Ca–Zn alloy, J. Mater. Sci. 25 (1990) 4579–4586. [5] N.A. Belov, E.A. Naumova, V.D. Ilyukhin, V.V. Doroshenko, Structure and mechanical properties of Al – 6% Ca – 1% Fe alloy foundry goods, obtained by die casting, Tsvetnye Met. 3 (2017) 69–75 https://doi.org/10.17580/tsm.2017.03.11. [6] N.A. Belov, E.A. Naumova, T.K. Akopyan, Effect of calcium on structure, phase composition and hardening of Al–Zn–Mg alloys containing up to 12 wt.%Zn, Mater. Res. 18 (2015) 1384–1391 http://dx.doi.org/10.1590/1516-1439.036415. [7] N.A. Belov, K.A. Batyshev, V.V. Doroshenko, Microstructure and phase composition of the eutectic Al – Ca alloy, additionally alloyed with small additives of zirconium, scandium and manganese, Nonferrous Met. 2 (2017) 49–54 http://dx.doi.org/10. 17580/nfm.2017.02.09. [8] N.A. Belov, E.A. Naumova, A.N. Alabin, I.A. Matveeva, Effect of scandium on structure and hardening of Al–Ca eutectic alloys, J. Alloy. Comp. 646 (2015) 741–747 https://doi.org/10.1016/j.jallcom.2015.05.155. [9] D.G. Altenpohl, Aluminum: Technology, Applications, and Environment – a Profile of a Modern Metal, sixth ed., The Aluminum Association, Inc., Washington, D.C., 1998. [10] S.A. Nikulin, A.B. Rozhnov, S.O. Rogachev, V.M. Khatkevich, V.A. Turchenko, E.S. Khotulev, Investigation of structure, phase composition, and mechanical properties of Zr-2.5% Nb alloy after ECAP, Mater. Lett. 169 (2016) 223–226 https://doi.org/10.1016/j.matlet.2016.01.148. [11] R.Z. Valiev, Y. Estrin, Z. Horita, T.G. Langdon, M.J. Zehetbauer, Y.T. Zhu, Fundamentals of superior properties in bulk nanoSPD materials, Mater. Res. Lett. 4 (2016) 1–21 https://doi.org/10.1080/21663831.2015.1060543. [12] Y. Estrin, M.J. Zehetbauer, Niche applications of bulk nanostructured materials processed by severe plastic deformation, in: Michael J. Zehetbauer, Yuntian Theodore Zhu (Eds.), Bulk Nanostructured Materials, Wiley-VCH, Weinheim, Germany, 2009, pp. 635–648 https://doi.org/10.1002/9783527626892.ch28. [13] F.Y. Dong, P. Zhang, J.C. Pang, Y.B. Ren, K. Yang, Z.F. Zhang, Strength, damage and fracture behaviors of high-nitrogen austenitic stainless steel processed by highpressure torsion, Scr. Mater. 96 (2015) 5–8 https://doi.org/10.1016/j.scriptamat. 2014.09.016. [14] A.V. Polyakov, D.V. Gunderov, G.I. Raab, Evolution of microstructure and mechanical properties of titanium Grade 4 with the Increase of the ECAP-Conform passes, Mater. Sci. Forum 667–669 (2011) 1165–1170 https://doi.org/10.4028/ www.scientific.net/MSF.667-669.1165. [15] J. Estrin, M. Murashkin, R. Valiev, Ultrafine-grained aluminium alloys: processes, structural features and properties, in: R.N. Lumley (Ed.), Fundamentals of Aluminium Metallurgy, Woodhead Publishing, UK, 2011, pp. 468–503 https://doi. org/10.1533/9780857090256.2.468. [16] D. Orlov, Y. Beygelzimer, S. Synkov, V. Varyukhin, N. Tsuji, Z. Horita, Plastic flow, structure and mechanical properties in pure Al deformed by twist extrusion, Mater.
7