Effect of Sc, Hf, and Yb Additions on Superplasticity of a Fine-Grained Al-0.4%Zr Alloy

Abstract

This research was undertaken to study the way deformation behaves in ultrafine-grained (UFG)-conducting Al-Zr alloys doped with Sc, Hf, and Yb. All in all, eight alloys were studied with zirconium partially replaced by Sc, Hf, and/or Yb. Doping elements (X = Zr, Sc, Hf, Yb) in the alloys totaled 0.4 wt.%. The choice of doping elements was conditioned by the possible precipitation of Al₃X particles with L1₂ structure in the course of annealing these alloys. Such particles provide higher thermal stability of a nonequilibrium UFG microstructure. Initial coarse-grained samples were obtained by induction casting. A UFG microstructure in the alloys was formed by equal-channel angular pressing (ECAP) at 225 °C. Superplasticity tests were carried out at temperatures ranging from 300 to 500 °C and strain rates varying between 3.3 × 10⁻⁴ and 3.3 × 10⁻¹ s⁻¹. The highest values of elongation to failure are observed in Sc-doped alloys. A UFG Al-0.2%Zr-0.1%Sc-0.1%Hf alloy has maximum ductility: at 450 °C and a strain rate of 3.3 × 10⁻³ s⁻¹, relative elongation to failure reaches 765%. At the onset of superplasticity, stress (σ)–strain (ε) curves are characterized by a stage of homogeneous (uniform) strain and a long stage of localized plastic flow. The dependence of homogeneous (uniform) strain (ε_eq) on test temperature in UFG Sc-doped alloys is increasing uniformly, which is not the case for other UFG alloys, with ε_eq(T) dependence peaking at 350–400 °C. The strain rate sensitivity coefficient of flow stress m is small and does not exceed 0.26–0.3 at 400–500 °C. In UFG alloys containing no Sc, the m coefficient is observed to go down to 0.12–0.18 at 500 °C. It has been suggested that lower m values are driven by intensive grain growth and pore formation in large Al₃X particles, which develop specifically at an ingot crystallization stage.

1. Introduction

Avionic systems of modern aircraft are manufactured using thin composite wires made of low-doped aluminum alloys [1,2]. Such a modern aircraft onboard network is several kilometers long, which is why the replacement of copper wires with aluminum ones helps reduce the weight of an onboard network and improve the efficiency of aviation technology. Thin aluminum wires find applications in the automotive industry [3] and in the modern power industry [4]. Eutectic aluminum alloys with a high content of rare-earth elements (REEs) and commercial conductor Al-Mg-Si alloys are often used to produce wires [1,2,5,6]. The thermal stability of such commercial Al-REEs alloys and conventional Al-Mg-Si alloys is insufficient; therefore, many research teams are now actively engaged in designing novel heat-resistant, high-strength conductor aluminum Al-REEs alloys [6,7,8,9,10,11,12,13,14].

Specific requirements as to microstructure thermal stability and properties during long-term (several thousand hours) operation at 200–250 °C are imposed on thin aviation wires [1]. In this context, novel conductor dispersoid-hardened Al alloys are being actively developed. Scandium is one of the most effective doping elements for aluminum [15,16,17,18,19,20,21,22,23], but high prices for Al-Sc master alloys hold back its extensive use. In this regard, researchers are developing multicomponent conductor aluminum alloys with improved strength, electrical conductivity, and thermal stability ensured by cheaper REEs and transition metals (TM) (see [15,17], etc.).

The Al-Zr system is one of the most popular for conductor aluminum alloys [24,25,26,27,28,29,30,31], but disadvantages of zirconium include low diffusibility of Zr atoms in aluminum [32,33,34] and discontinuous precipitation of a solid solution during annealing of Al-Zr alloys [35,36,37,38,39]. Therefore, Al-Zr alloys are often additionally doped with some other elements that accelerate the diffusibility of Zr atoms. Most often, Al-Zr alloys are doped with X = Er, Hf, Yb, Y, Si, Ti, W, etc. [40,41,42,43,44,45,46,47,48,49,50,51,52,53,54,55,56,57,58,59,60,61,62,63,64,65,66]. This allows for the precipitation of Al₃(Zr,X) particles of variable composition at lower temperatures compared to precipitation temperatures of Al₃Zr particles. The wires are made by deformation (rolling, drawing, etc.), and precipitation of Al₃(Zr,X) nanoparticles at low temperatures ensures better thermal stability of an ultrafine-grained (UFG) microstructure of a deformed aluminum alloy [67,68].

Currently, researchers are focused on studying the mechanical properties and electrical conductivity of conductor aluminum alloys at room temperature. The deformation behavior of conductor aluminum alloys at elevated temperatures remains under researched, although the deformation of an aluminum workpiece in most cases takes place at elevated temperatures. There are many scientific papers on the superplasticity of fine-grained conductor 6XXX Al alloys [69,70,71,72,73,74,75,76,77,78], but only few works on Mg-free UFG Al-1%Zr alloys [79] and Al-0.5%Mg-Sc alloys [80,81]. Optimal temperatures and strain rates will help to produce aluminum wires with a minimum number of breaks. An optimized composition of conducting aluminum alloys can bring down the number of breaks by reducing the number and size of primary Al₃Zr particles formed during the crystallization of Al-Zr-X alloys.

This study focuses on Al-Zr-Sc, Al-Zr-Hf, and Al-Zr-Yb alloys, as well as alloys simultaneously alloyed with several doping elements. Precipitation of Al₃(Zr,Sc) and Al₃(Zr,Hf) secondary particles in Al-Zr-(Sc,Hf) alloys is still more intensive at lower temperatures compared to Al₃Zr in Al-Zr alloys [15,40,43,45,68]. Therefore, during low-temperature heating, it is possible to ensure precipitation of a larger volume fraction of Al₃(Zr,X) secondary particles compared to that of Al₃Zr particles. The low solubility of ytterbium (Yb) in molten aluminum at ordinary crystallization temperatures [82,83,84,85] and low solubility of Yb in Al at room and medium temperature [86,87] helps to obtain alloys with precipitated fine primary Al₃Yb particles [88]. Of note, secondary precipitation of Al₃(Yb,Zr) and Al₃(Yb,Sc) secondary particles is still more intensive at lower temperatures compared to Al₃Zr and Al₃Sc precipitates [41,56,89,90,91,92,93]. Thus, it is possible to make a fine-grained aluminum microstructure more stable at low annealing temperatures, as well as to further refine an ingot macrostructure. It should be noted that a large volume fraction of strong primary particles and their large size can lead to wire breaks during low-temperature rolling and drawing. Therefore, in this research, we limited the concentration of doping elements (no more than 0.4 wt.%), which was designed based on the results of previous studies.

The high stability of a UFG microstructure in aluminum alloys at lower heating temperatures allows for stable properties of thin wires operating at 200–250 °C [7,8,43,67,68]. Moreover, the stability of a nonequilibrium UFG microstructure at low temperatures will enable the low-temperature superplasticity effect in UFG aluminum alloys [69,70,71,72,73,74,75,76,80,81]. Raising the ductility of aluminum alloys while reducing their optimal deformation temperatures can help to increase the energy efficiency of production, in particular to reduce the number of breaks during the manufacturing of thin wires and to reduce the wear of equipment used in wire manufacturing. A drop in deformation temperatures and a rise in the rate of manufacture of thin bimetallic wires helps to further reduce the intensity of copper diffusion into the surface of an aluminum wire and, as a result, to reduce the intensity of brittle intermetallic compounds formation at the Al/Cu interface [68,94,95,96].

This study aims to research the deformation behavior of UFG conductor aluminum alloys at elevated temperatures.

2. Materials and Methods

Aluminum alloys with the content of doping elements totaling 0.4 wt.%, which corresponds to the solubility limit of zirconium in aluminum in selected induction casting modes, are investigated in this study. The chemical composition of the alloys under investigation is presented in

Table 1. Aluminum alloys: chemical composition and casting modes.

Alloy No.	Doping Elements, wt.%/at.% 1					Casting Modes
	Zr	Sc	Hf	Yb	∑
1	0.4	-	-	-	0.4	Argon purging before melting: 3 cycles Argon purging during heating: 3 cycles Induction stirring Cooling: 50–250 s, under vibration Time to melt: 6–10 min Melt temperature: 800 °C Holding time before pouring: 60 min Pouring temperature: 760 °C
	0.118	-	-	-	0.118
2	0.3	0.1	-	-	0.4
	0.089	0.167	-	-	0.255
3	0.3	-	0.1	-	0.4
	0.089	-	0.015	-	0.104
4	0.3	-	-	0.1	0.4
	0.089	-	-	0.016	0.105
5	0.2	0.1	0.1	-	0.4
	0.059	0.167	0.015	-	0.241
6	0.2	-	0.1	0.1	0.4
	0.059	-	0.015	0.016	0.090
7	0.2	0.1	-	0.1	0.4
	0.059	0.167	-	0.016	0.242
8	0.1	0.1	0.1	0.1	0.4
	0.030	0.167	0.015	0.016	0.228

¹—numerator–concentration of doping elements (wt.%), denominator–calculated concentration of doping elements (at.%).

.

Aluminum alloy workpieces of 20 × 20 × 160 mm were obtained by vacuum induction casting in the modes indicated in Table 1. An Indutherm VTC 200 V vacuum casting machine (Indutherm GmbH, Walzbachtal, Germany) was used to produce ingots that were cast in a copper mold of 22 × 22 × 160 mm. A 150 cm³ ceramic crucible was used. Alloys were made from aluminum A99, and master alloys such as Al-10%Zr, Al-2%Sc, Al-2.5%Yb, and Al-3%Hf were obtained by induction casting, followed by rolling into a foil 0.2 mm thick.

Equal-channel angular pressing (ECAP) was used to form a UFG microstructure in aluminum alloys. A detailed description of this technology can be found in [97,98]. ECAP helps to form a homogeneous UFG microstructure in aluminum alloy ingots, the parameters of which are close to those of a microstructure of a thin aluminum wire (see [67,68]). The above workpieces were subjected to 4 ECAP cycles to form a UFG microstructure. ECAP was performed at 250 °C using a Ficep HF400L hydraulic press (Ficep^® S.P.A., Varese, Italy). Pressing was arranged using square-section tooling, with a channel intersection angle of 90°. ECAP was performed without billet rotations at a strain rate of 0.4 mm/s.

The resulting flat dogbone-shaped specimens with a working part of 2 × 2 × 3 mm were subjected to mechanical tests. Tensile tests were conducted using a Tinius Olsen H25K-S machine (Tinius Olsen Ltd., Surrey, UK) at a strain rate of 3.3 × 10⁻³ s⁻¹ (tensile rate of 10⁻² mm/s). Tests were run at temperatures ranging from 20 to 500 °C. During tests, a stress (σ)–strain (ε) curve was registered and used to determine ultimate tensile strength (σ_b) and elongation to failure (δ). Three specimens were tested at the above temperatures and strain rates.

An alloy microstructure was studied with a Leica DM IRM metallographic microscope (Leica Microsystems GmbH, Wetzlar, Germany), a Jeol JSM-6490 scanning electron microscope (SEM, Jeol Ltd., Tokyo, Japan), and a Jeol JEM-2100F transmission electron microscope (TEM, Jeol Ltd., Tokyo, Japan). Specimens were preliminarily subjected to mechanical grinding and polishing, as well as electropolishing (3 A, 30 V) for 1 min in CrO₃ + H₃PO₄ electrolyte. An alloy microstructure was exposed by etching in a solution of HF (15 mL) + HNO₃ (10 mL) + glycerol (35 mL). The fractographic analysis that followed tensile tests was performed using a Jeol JSM-6490 SEM. The microstructure of specimens after tensile tests was studied using a Leica IM DRM metallographic microscope. Microhardness (HV) was measured using a Qness A60+ hardness tester (Qness GmbH, Golling, Austria) under a 50 g load. The average accuracy of HV measurement was ±15 MPa.

3. Results

3.1. Initial State Characterization

Cast alloys have a homogeneous macrostructure with large columnar crystals. The length of columnar crystals in the central part of thin sections may reach several millimeters. Along the edges of thin sections, there is an area of equiaxed grains, the average size of which is 10–50 μm, which leads to a fairly large spread of microhardness values registered in the center and along the edge of a thin section. No significant differences in macrostructures of alloy specimen Nos. 1–8 were observed.

The microstructure of cast and UFG alloys contains primary particles, the size and composition of which depend on the chemical composition of an alloy. Alloys doped with 0.3 and 0.4% Zr have single-micron Al₃Zr particles. UFG alloy Nos. 4, 6, 7, and 8 after ECAP have a large amount of light spherical and elongated flat Al₃Yb particles (Figure 1c–f). Single particles of a complex composition containing Fe, Si, and Ni can be observed in all the alloys; some particles have oxygen in their composition (Figure 1).

The microstructure of UFG alloys is typical of severely deformed metals; the average fragment size is close to 0.5 μm (Figure 2a). No significant differences in the microstructure parameters of UFG alloy Nos. 1–8 were observed. No abnormally large grains were found in the microstructure of UFG alloys. A dark-field TEM image shows primary Al₃(Zr,X) nanoparticles that were breaking down at an ingot crystallization stage (Figure 2b).

Table 2. Mechanical properties of UFG alloys at room temperature.

Alloy No.	σb, MPa	δ, % 1	HV, MPa
1	115	15–20	410 ± 10
2	140	20–25	455 ± 15
3	115	20–25	385 ± 15
4	125	15–20	410 ± 10
5	145	25–30	465 ± 10
6	125	20–25	415 ± 10
7	135	20–25	480 ± 20
8	140	20–25	455 ± 10

¹—Relative elongation to failure was determined by changes in the length of a gage zone of a specimen after tensile tests.

presents the results of studying the mechanical properties of UFG alloy specimens. σ(ε) curves are typical of highly deformed materials; strain hardening is almost immediately followed by strain localization. Maximum ultimate tensile strength values (σ_b = 140–145 MPa) are observed in UFG alloy Nos. 2, 5, and 8. Relative elongation to failure in most UFG alloys is 15–20% (see Table 2).

Figure 3 shows the results of the fractographic analysis of UFG alloy No. 5 specimens. The results of the fractographic analysis of the remaining specimens are presented in Appendix A. According to the classification provided in [99], fractures that occur in specimens after tensile tests at room temperature are characterized by ductility (Figure 3a, see Appendix A). Specimens during tensile tests are destroyed to form a localization zone (see photographs of specimens in the upper left corner of Figure 3a). The central zone of fractures has dimples of various sizes (Figure 3b). UFG alloys characterized by increased strength and high elongation to failure have smaller dimples as compared to other UFG alloys (see Appendix A).

3.2. Tests at Elevated Temperatures

Figure 4 and Appendix B show the stress (σ)–strain (ε) curves in UFG alloys at elevated tensile test temperatures. The σ(ε) tensile curves at low tensile temperatures (300–350 °C) are characterized by a pronounced stage of strain hardening, followed by a stage of localized plastic deformation. A uniform plastic flow stage is very short (Figure 4). With an increase in test temperatures, flow stress decreases, and elongation to failure in UFG alloys increases (

Table 3. Tensile tests of UFG alloys at elevated temperatures and a strain rate of 3.3 × 10⁻³ s⁻¹.

Alloy No.	300		350		400		450		500
	σb, MPa	δ, %	σb, MPa	δ, %	σb, MPa	δ, %	σb, MPa	δ, %	σb, MPa	δ, %
1	47 ± 2	185 ± 10	20 ± 2	240 ± 20	14 ± 1	230 ± 20	10 ± 0.5	215 ± 20	8 ± 0.5	155 ± 15
2	53 ± 4	195 ± 10	20 ± 1	255 ± 20	11 ± 0.5	475 ± 25	7 ± 0.5	400 ± 30	5 ± 0.3	685 ± 40
3	28 ± 2	190 ± 10	23 ± 1	180 ± 15	15 ± 1	215 ± 15	9 ± 0.5	145 ± 15	7 ± 0.5	150 ± 15
4	31 ± 2	190 ± 10	27 ± 2	185 ± 15	17 ± 1	185 ± 15	9 ± 0.5	135 ± 10	9 ± 0.5	135 ± 10
5	31 ± 2	220 ± 15	21 ± 1	280 ± 25	12 ± 0.5	325 ± 25	8 ± 0.5	765 ± 40	6 ± 0.5	525 ± 40
6	38 ± 3	185 ± 10	26 ± 2	175 ± 15	16 ± 1	145 ± 10	13 ± 1	155 ± 15	7 ± 0.5	155 ± 15
7	30 ± 1	210 ± 20	20 ± 1	295 ± 25	11 ± 1	455 ± 30	8 ± 0.5	360 ± 35	9 ± 0.5	250 ± 20
8	31 ± 2	260 ± 20	21 ± 2	265 ± 15	12 ± 1	325 ± 25	8 ± 0.5	385 ± 30	7 ± 0.5	320 ± 30

). At test temperatures of 400–500 °C, the σ(ε) curves for UFG Sc-doped alloys have all the signs of superplastic flow: flow stress is very low, while elongation to failure exceeds several hundred percent (Table 3, Figure 4b). The σ(ε) curves for UFG Sc-doped alloy Nos. 2, 5, 7, and 8 show long stages of uniform plastic flow, followed by a long stage of localized plastic flow (Figure 4b, Appendix B).

At a localized plastic deformation stage, necking is observed in the gage zone of a specimen. Several localized deformation zones occur in Sc-doped alloys at 500 °C (Figure 5). Intense formation of localized deformation zones is generally observed at elevated temperatures and low strain rates.

The results of tensile tests of UFG alloys at elevated temperatures are presented in Table 3.

Data in Table 3 show that elongation to failure nonmonotonically (with a peak) depends on test temperatures; flow stress decreases monotonically along with an increase in test temperatures. The highest superplasticity is registered in Sc-doped alloy Nos. 2, 5, 7, and 8 (Table 3). At 400 °C and a strain rate of 3.3 × 10⁻³ s⁻¹, maximum elongation to failure (δ = 455–475%) is observed in UFG alloy Nos. 2 and 7 (Table 3). UFG alloy No. 5 has maximum ductility at elevated temperatures: at 450 °C and a strain rate of 3.3 × 10⁻³ s⁻¹, elongation to failure reaches 765% (Table 3). UFG alloy No. 2 also has high superplastic characteristics: at 500 °C and a strain rate of 3.3 × 10⁻³ s⁻¹, elongation to failure reaches 685%.

Figure 6 shows the dependence of uniform strain (ε_eq) on test temperature in UFG alloys. When a long stage of uniform plastic flow appeared on the σ(ε) curve, ε_eq was taken as deformation corresponding to an early stage of strain localization (onset of stress reduction).

It is noteworthy that based on the nature of ε_eq(T) dependence, all UFG alloys can be divided into two large groups. Alloy No. 1, 3, 4, and 6 with low ductility at elevated temperatures (400–500 °C) (Table 3) display a nonmonotonic ε_eq(T) dependence. Figure 6 shows that ε_eq(T) dependence in UFG alloy Nos. 1, 3, 4, and 6 reaches its maximum at 350–400 °C. With a further increase in test temperatures, ε_eq decreases. It stands to mention that maximum uniform strain values for UFG alloy Nos. 1, 3, 4, and 6 are small and do not exceed 70–90% (Figure 6). ε_eq(T) dependence in UFG alloy Nos. 2, 5, 7, and 8 is different: with an increase in test temperatures, there is a monotonic increase in ε_eq. UFG alloy No. 2 is an exception since it demonstrates a slight decrease in ε_eq at 500 °C (Figure 6), but the scale of this decrease only slightly exceeds ε_eq error.

The effect of strain rate on the results of superplasticity tests was studied at 400 and 500 °C. The test results are presented in

Table 4. Results of tensile tests of UFG alloy specimens at elevated temperatures and strain rates. Numerator–flow stress (MPa), denominator–elongation to failure (%).

Alloy No.	T = 400 °C				T = 500 °C
	3.3 × 10−4	3.3 × 10−3	3.3 × 10−2	3.3 × 10−1	3.3 × 10−4	3.3 × 10−3	3.3 × 10−2	3.3 × 10−1
1	-	14 ± 0.5	25 ± 1	-	-	8 ± 0.5	11 ± 0.5	14 ± 1
	-	230 ± 20	130 ± 10	-	-	155 ± 10	160 ± 15	180 ± 15
2	7 ± 0.5	11 ± 0.5	26 ± 1	-	3 ± 0.2	5 ± 0.3	13 ± 1	-
	335 ± 30	475 ± 35	280 ± 20	-	435 ± 35	685 ± 35	385 ± 25	-
3	-	15 ± 1	34 ± 2	45 ± 2	-	7 ± 0.4	9 ± 0.5	13 ± 1
	-	215 ± 20	155 ± 15	105 ± 10	-	150 ± 10	200 ± 15	195 ± 15
4	-	17 ± 1	34 ± 2	47 ± 3	-	9 ± 0.5	11 ± 0.5	14 ± 1
	-	185 ± 15	150 ± 15	105 ± 10	-	135 ± 10	110 ± 10	150 ± 15
5	7 ± 0.5	12 ± 0.5	27 ± 2	-	5 ± 0.3	6 ± 0.3	11 ± 1	-
	340 ± 30	325 ± 25	235 ± 20	-	500 ± 30	525 ± 35	515 ± 35	-
6	-	16 ± 1	29 ± 2	35 ± 2	-	7 ± 0.4	10 ± 0.5	14 ± 1
	-	145 ± 15	130 ± 10	115 ± 10	-	155 ± 15	110 ± 10	150 ± 10
7	-	11 ± 0.5	24 ± 1	44 ± 3	-	9 ± 0.5	12 ± 1	29 ± 2
	-	455 ± 40	275 ± 30	160 ± 15	-	250 ± 25	440 ± 35	215 ± 15
8	-	12 ± 1	21 ± 1	45 ± 3	-	7 ± 0.4	14 ± 1	30 ± 2
	-	325 ± 35	280 ± 25	135 ± 10	-	320 ± 30	300 ± 25	145 ± 10

. The analysis of data in Table 4 shows that an increase in strain rates triggers an increase in flow stress, which is a common effect in the superplasticity tests of fine-grained materials [100]. Elongation to failure nonmonotonically (with a peak) depends on the test temperatures of UFG aluminum alloys (Table 4). In the region of high-strain-rate superplasticity at a strain rate of 3.3 × 10⁻¹ s⁻¹, elongation to failure in UFG aluminum alloys reaches 105–160% at 400 °C, while at 500 °C for some UFG alloys, δ is close to 200% (Table 4). At higher strain rates, uniform deformation (ε_eq) is small, and a localized deformation stage is mostly observed on σ(ε) curves (Figure 7). Figure 8 shows that the higher the strain rate, the shorter the stage of localized deformation and the higher the flow stress.

Figure 8 and Appendix B show the results of the fractographic analysis of UFG alloy specimens after superplasticity tests at 400 and 500 °C. Fractures in all specimens are ductile in nature and represent a set of dimples of various sizes. With an increase in deformation temperatures, dimples diminish in size along with a cross-sectional area of fractures (Figure 8). UFG alloy specimens with increased ductility have an extremely small cross-sectional area of fractures; fractures show single dimples of 10–20 µm. A small cross-sectional area of fractures indirectly indicates a high strain localization, which corresponds to data concerning little uniform strain (ε_eq) in these alloys (Figure 8). Such a conclusion is underpinned by shape analysis of specimens after tensile tests, the photos of which are shown in the upper right corner of Figure 8 and in Appendix B. It should be noted that single dimples are visible in fractures of some specimens tested at 400 °C, with the size of dimples reaching 50–100 μm (Appendix B). In our opinion, cavitation fracture is the reason for such defects.

Figure 9 presents the results of the microstructure investigations of the UFG aluminum alloy samples from series #5 after the superplasticity tests with the strain rate of 3.3 × 10⁻³ s⁻¹ at 450 °C and 500 °C. The samples of this series had the highest value of the elongation to failure (Table 2). The area of investigation is outlined by a yellow dashed line in Figure 9a.

The pores are seen on the sample surfaces after the tension tests, and the sizes and quantity of the pores increase near the destruction zone (Figure 9b,c). Note that according to the model [101], the presence of the secondary particles nucleating when heating the UFG aluminum alloy is the origin of the pore formation in the superplasticity conditions. The dislocation-disclination type defects forming around large secondary particles in the superplasticity conditions can lead to the acceleration of the cavitation destruction of the UFG alloys [102,103,104,105]. A similar effect was described earlier for the case of superplasticity of UFG conductor aluminum alloys Al-0.5%Mg-Sc [80,81].

Figure 9d,e presents the microstructure images of the UFG alloy samples near the destruction areas. This area is outlined by the yellow dashed line in Figure 9a. One can see in Figure 9d that the mean grain sizes in the destruction area in the samples after the superplasticity tests at 450–500 °C were ~10–15 μm. The mean grain sizes (d ~ 10–15 μm) exceeded the initial grain size in the UFG alloys (~0.5 μm) essentially. The result obtained evidences an intensive grain growth in the superplasticity conditions. Note that the grains had the shapes close to the equiaxial ones while the pores were located preferentially at the grain boundaries (Figure 9d,e).

4. Discussion

ECAP is an effective technology to form a UFG microstructure in aluminum alloys [97,98,106,107]. A distinct ECAP advantage is the opportunity to form a large fraction of high-angle grain boundaries (HAGBs) in a fine-grained material [98,106,107]. HAGBs in the microstructure of a fine-grained material help to launch the grain boundary sliding (GBS) effect, which, under optimal temperature–rate deformation conditions, is the main mechanism of the low-temperature superplastic deformation of fine-grained materials [100,108]. The third factor contributing to the higher superplastic characteristics of UFG alloys is the nonequilibrium state of their high-angle grain boundaries. As shown in [98,109,110,111,112], during ECAP, HAGBs accumulate dislocation defects, leading to an increase in the free volume of grain boundaries and a decrease in the activation energy of superplastic deformation. During the superplastic deformation of UFG alloys, their grain boundaries actively interact with lattice dislocations, thus triggering an increase in the free (excess) volume of high-angle grain boundaries [98,110,111,112]. Thus, ECAP provides all the necessary conditions conducive to the superplasticity effect in aluminum alloys. The superplasticity of UFG aluminum alloys produced by ECAP has been studied thoroughly in many works (see, for example, [80,81,106,113,114,115,116,117]); therefore, we will not dwell on their analysis. For the purposes of analyzing the results that we obtained, we should only emphasize that UFG-conducting aluminum alloys exhibit very good superplastic characteristics (Table 3 and Table 4).

Let us analyze superplastic deformation in UFG-conducting aluminum alloys. The main rheological superplastic flow equation usually runs as follows [100,108]:

$$\dot{\epsilon }=A{\left({\mathsf{\sigma }}^{*}/G\right)}^{1/m}{\left(b/d\right)}^p\left(D_{eff}/b^2\right)\left(G\Omega /kT\right)$$

(1)

where m is the strain rate sensitivity coefficient of flow stress, the value of which depends on strain rate $\dot{\epsilon }$ (m is usually 0.5 under optimal superplasticity conditions); p is a numerical parameter equal to 2 or 3 [109]; b is the Burgers vector; G is the shear modulus; k is the Boltzmann constant; σ^* is flow stress; and D_eff is the effective diffusion coefficient under superplasticity conditions. It is usually assumed that σ^* = σ_b, D_eff = D₀exp(Q_eff/kT), where Q_eff is the effective activation energy of superplastic flow. Q_eff is usually close to grain boundary diffusion activation energy Q_b [108], but in the case of UFG metals, it often turns out to be less than Q_b [98,110,111], and in the case of multiphase materials or in the case of diffusion creep, it can be greater than Q_b [100].

The rheological superplasticity equation can also be presented in such a way that contributions of intergranular (${\dot{\epsilon }}_b)$ and intragranular deformation (${\dot{\epsilon }}_v)$ are separated:

$$\dot{\epsilon }={\dot{\epsilon }}_b+{\dot{\epsilon }}_v$$

(2)

To describe the grain boundary sliding (GBS) rate ${\dot{\epsilon }}_b$, the following equation is usually used:

$${\dot{\epsilon }}_b=A_b{\left(\mathsf{\sigma }/G\right)}^2{\left(b/d\right)}^2\left(G\Omega /kT\right)\left(\delta D_b/b^3\right)$$

(3)

where A_b is a numerical coefficient equal to ~100 [98,108]. To calculate intragranular strain rate ${\dot{\epsilon }}_v$, the usual equation for strain rate under power-law creep [100,108] is used.

The strain rate sensitivity coefficient of flow stress $m=\ln \left({\mathsf{\sigma }}_b\right)/\ln \left(\dot{\epsilon }\right)$ can be determined by the slope of flow stress–strain rate dependence in logarithmic coordinates. These dependencies in the UFG aluminum alloys under study are shown in Figure 10. The charts provided herein show that rather low m values that do not exceed 0.26–0.30 are observed in all alloys at 400 and 500 °C (Figure 10). Such results prove that the selected temperature–rate modes of UFG alloy deformation are not optimal.

It is worth noting that the values of the m coefficient in most UFG alloys at 500 °C are somewhat lower than at 400 °C. The m values for Sc-doped alloys turn out to be close to each other at 400 and 500 °C. This is rather unexpected, since higher elongations to failure are observed at elevated temperatures (Table 3 and Table 4). We reckon that there are two main reasons why the superplastic deformation modes of UFG-conducting aluminum alloys are suboptimal at elevated temperatures.

First, it should be pointed out that dynamic grain growth is observed in UFG aluminum alloys at elevated temperatures and strain rates (see Figure 9). According to Equations (1) and (3), this leads to a decrease in optimal superplastic strain rates, as well as to changes in strain hardening and an increase in flow stress [112,113,114,115,116,117,118,119,120]. Moreover, it must be mentioned that intensive dynamic grain growth leading to the formation of a coarse-grained microstructure is one of the main reasons for the higher strain hardening and lower elongation to failure of UFG materials under deformation at elevated temperatures.

It should be emphasized that UFG Sc-doped alloys exhibit the highest superplastic characteristics [80,81,114,115,116,117,118]. Al₃Sc particles and Sc-containing intermetallic compounds of variable composition with L1₂ structure are among the most effective stabilizers for a nonequilibrium UFG microstructure of aluminum alloys [18,19,20,21,121,122,123,124]. This suggests that UFG Sc-doped alloys have the lowest rate of dynamic grain growth and, as a consequence, the smallest grain size under superplasticity conditions. In accordance with (1), this allows for higher optimal strain rates and maximum superplastic characteristics of UFG alloys.

The second factor is secondary particles present in a UFG alloy microstructure, including relatively large primary particles formed at an ingot crystallization stage. Large strong particles trigger the acceleration of pore nucleation under superplastic deformation conditions [101] and consequently premature cavitation fracture in UFG alloys. According to a model provided in [101], pore nucleation occurs because dislocation glide is hugely impeded, and dislocations and disclinations emerge around particles. Cavitation fractures in UFG-conducting alloys under superplasticity conditions tally with the results of fractographic analysis (Figure 8, Appendix C). As noted above, one of the reasons for large pores on the fracture surfaces in UFG alloy specimens can be large particles in the failure zone. This conclusion indirectly agrees with the results of superplasticity tests. As can be seen from Table 3 and Table 4, maximum elongation to failure is observed in UFG alloys with reduced Zr content. Large fan-shaped particles of Al₃Zr and Al₃Sc, which are often formed by discontinuous precipitation [33,34,35,36,37,38,39,80,81,124,125,126,127], can be one of the main factors for accelerated cavitation fracture in UFG alloys under superplasticity conditions. Cavitation fracture accompanying superplasticity in UFG-conducting alloys with elevated Sc content was described in [80,81].

In conclusion, let us analyze the way the type and composition of secondary particles affect the superplasticity characteristics of UFG aluminum alloys. As shown above, the highest superplasticity characteristics are achieved in UFG Sc-doped alloys (Figure 11).

When analyzing the results obtained, it should be borne in mind that Zr, Sc, Hf, and Yb have significantly different distribution coefficients in aluminum. The way doping elements are distributed in the alloy during crystallization can be described using distribution coefficient K = C_S/C_L, where C_S and C_L stand for concentrations of doping elements in the liquid and solid phases, respectively [128]. At K < 1, doping elements concentrate in the liquid phase during crystallization and form grain boundary segregations after crystallization. At K > 1, doping elements tend to concentrate in the grain volume during crystallization. At K = 1, doping elements are distributed evenly between grain boundaries (dendritic boundaries) and the volume of grains (dendritic volume). In aluminum, zirconium (K_Zr = 2.5–2.54 [128]) and hafnium (K_Hf = 4.0–6.01 [128]) are horophobic elements and, after crystallization, are predominantly found in the aluminum crystal lattice. Ytterbium in aluminum (K_Yb = 0.08 [128]) is a horophilic element and forms segregations along grain boundaries or dendritic boundaries. Scandium is uniformly distributed in aluminum (K_Sc = 0.9–1.0 [128]). Differences in the way doping elements are distributed can lead to differences in the way particles precipitate during the heating of UFG aluminum alloys.

The second factor that must be taken into account when analyzing the results is that due to differences in the atomic masses of Al (26.98 a.m.u.), Zr (91.224 a.m.u.), Sc (44.956 a.m.u.), Hf (178.49 a.m.u.), and Yb (173.044 a.m.u.), the concentration of Sc in at.% in the alloys under study significantly exceeds that of all other doping elements (Zr, Hf, and Yb) (Table 1). Therefore, in Sc-doped alloy Nos. 2, 5, 7, and 8, the volume fraction of precipitating (f_v) Al₃X particles will be noticeably higher than in other alloys (No. 1, 3, 4, and 6). Table 1 shows that alloy Nos. 2, 5, and 7 are expected to have the largest volume fraction of precipitated particles. A nonequilibrium UFG microstructure in Sc-doped alloys will stabilize better due to the Zener relation: d_Z = α₁f_v/R, where d_Z is the stable grain size, R is the particle size, and α₁ is a numerical coefficient depending on the particle shape [129]. The reduced ductility of Sc-doped alloy No. 7 compared with alloy Nos. 2 and 5 (Table 3 and Table 4) stems from the fact that a big part of the total concentration of doping elements is made up of ytterbium atoms located along grain boundaries in the form of large primary Al₃Yb particles (Figure 2) and have little effect on stabilization of a UFG microstructure in the aluminum alloy. Alloys 3, 4, and 6 (Figure 11) with the lowest total concentration of doping elements display the lowest superplasticity (minimum elongation to failure) (Table 1), and therefore, we can expect the smallest volume fraction of secondary particles.

It should be noted that in accordance with the data presented in Table 1, one would expect increased superplastic characteristics in UFG alloy No. 2, which has the highest total concentration of doping elements (C_∑ = 0.241%). Figure 11 shows that at a deformation temperature of 450 °C, the highest values of elongation to failure are observed in UFG alloy No. 5 (C_∑ = 0.255%). Elongation to failure in UFG alloy No. 2 at 450 °C is lower than in UFG alloy No. 5. We reckon that the reason for this effect lies in the difference in the chemical composition and structure of precipitated particles in alloy Nos. 2 and 5.

As shown in [130,131,132], precipitating secondary Al₃(Zr,Sc) particles in Al-Zr-Sc alloys have the following structure: Al₃Sc core–Al₃Zr shell. The diffusion coefficient of scandium in aluminum [133,134,135,136] is noticeably higher than the diffusion coefficient of zirconium in aluminum [32]. As a result, Al₃Sc particles precipitate at lower temperatures or shorter heating times compared to Al₃Zr particles. At higher temperatures or longer heating times, a stable Al₃Zr shell is formed on the surface of precipitated Al₃Sc nanoparticles. An elevated Sc concentration in alloy No. 2 will trigger the formation of Al₃Sc particles in the crystal lattice and along the grain boundaries of the aluminum alloy.

Doping Al-Zr alloys with hafnium is also known to accelerate the formation of secondary Al₃Zr particles, but higher temperatures are required [40,43,68]. In our opinion, doping Al-Zr-Sc alloy with Hf during heating leads to secondary Al₃(Hf,Zr) particles that are formed at higher temperatures compared to Al₃(Sc,Zr) and Al₃Sc particles. Thus, it becomes possible to ensure the thermal stability of a fine-grained microstructure of Al-0.2%Zr-0.1%Sc-0.1%Hf alloy in a wide range of temperatures and strain rates. Small grain size in the alloys under study, in accordance with (1), leads to an increase in optimal rates of superplastic deformation and to an increase in the ductility of alloy No. 5.

5. Conclusions

1. This research focused on the superplastic characteristics of ultrafine-grained (UFG) conducting Al-Zr alloys doped with Sc, Hf, and Yb with a total content of doping elements not more than 0.4 wt.%. It was shown that the partial replacement of zirconium with rare-earth elements and other transition metals allows for higher ductility of UFG alloys at elevated temperatures. UFG alloys containing 0.1% Sc exhibit the highest superplasticity. A UFG Al-0.2%Zr-0.1%Sc-0.1%Hf alloy has maximum ductility: at 450 °C and a strain rate of 3.3 × 10⁻³ s⁻¹, its relative elongation to failure reaches 765%. A UFG Al-0.3%Zr-0.1%Sc alloy also has high superplastic characteristics: at 500 °C and a strain rate of 3.3 × 10⁻³ s⁻¹, its relative elongation to failure reaches 685%.

2. It was shown that at a deformation temperature of 400 °C, the strain rate sensitivity coefficient of flow stress (m) values are 0.26–0.28. An increase in deformation temperatures to 500 °C leads to a decrease in m values. The reason for low m values is dynamic grain growth and pore nucleation in large Al₃(Zr,X) particles.