**Casting and Solidification of Light Alloys**

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**Alexander Vorozhtsov**

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### **Contents**

Reprinted from: *Metals* **2020**, *10*, 772, doi:10.3390/met10060772 ................... **165**

### **About the Editor**

**Alexander Vorozhtsov**, Doctor of Physical and Mathematical Sciences, Professor, Head of High Energetic and Special Materials Laboratory of Tomsk State University, Russian Federation.

### **Preface to "Casting and Solidification of Light Alloys"**

The processes of casting and crystallization define the development of the final properties of light alloys, and optimization of these processes allows for improving efficiency of the use and applications of light alloys. The analysis and design of such processes requires special attention. The aim of this book is to present the latest achievements in technology, including advanced and modern approaches, structure development, and properties of the light alloys. Included are investigations of the effect of casting and crystallization on the structure and properties of the resulting light alloys (based on aluminum and magnesium to begin with) and, in particular, research connected with detailed analysis of the microstructure of light alloys obtained using various external influences of ultrasonic, vibration, magnetic, and mechanical processing on the casting and crystallization. The use of modern methods to study the properties of alloys in order to assess the effect of structure on the mechanical and functional properties of light alloys is also covered as are studies on the introduction of additives (modifiers, reinforcers, including nanosized ones, etc.) into the melt during the crystallization process, the technological properties of casting (fluidity, segregation, shrinkage, etc.), the structure and physicomechanical properties of light alloys. Finally, this book includes papers that focus on the relations of physicomechanical properties with the defective structure of light alloys and mathematical modeling of plastic deformation of dispersion-strengthening materials.

> **Alexander Vorozhtsov** *Editor*

### *Editorial* **Casting and Solidification of Light Alloys**

### **Alexander Vorozhtsov**

Laboratory of High Energetic and Special Materials, National Research Tomsk State University, Lenin av. 36, 634050 Tomsk, Russia<EMAIL_ADDRESS>Tel.: +7-3822-526-120

Received: 9 October 2020; Accepted: 15 October 2020; Published: 23 October 2020

### **1. Introduction and Scope**

At present, light alloys based on aluminum, magnesium, titanium, etc. have been intensively studied. Improving the properties of these alloys by modifying the structure and reinforcement makes it possible to solve the technological problem of increasing the working properties of mechanical engineering products while simultaneously reducing their weight. Concurrent to the study and development of methods for strengthening light alloys, casting technologies also require improvement since traditional casting technologies do not allow for the improvement of alloy properties when hardening particles are introduced into the liquid phase of the base metal. Modern methods of processing light alloys during casting make it possible to increase the physical and, ultimately, the operational properties of the obtained alloys by degassing, reducing the average grain size, increasing the uniformity of the alloy composition, reducing the amount of agglomerates and impurities at grain boundaries, increasing wettability, etc. The development of unconventional casting methods and the study of properties of light alloys are reported in the articles of this special issue.

### **2. Contributions**

This special issue consists of one review and 11 original research articles. Of these 12 articles, 10 articles are devoted to aluminum alloys, one article is devoted to titanium alloys and the review article is devoted to aluminum, titanium, copper and magnesium alloys.

The review article [1] provides an overview of recent advances in light alloys (aluminum, copper, titanium and magnesium alloys) modified using friction stir processing (FSP). The general mechanisms of the formation of subsurface gradient structures in metal alloys processed by FSP under various conditions are described. It is shown that FSP can be used to produce light alloys with subsurface gradient structure, composite subsurface gradient structure and "in-situ" composites.

In the first research article [2], aluminum alloys of the Al-Mg system with titanium diboride particles of different dispersion were obtained. The introduction of titanium diboride particles using ultrasonic treatment of the melt made it possible to significantly reduce the average grain size of the alloy AA5056. The greatest effect of structure refinement was obtained using a master alloy containing titanium diboride particles with a size of 1 μm. It was also found that the introduction of titanium diboride particles led to an increase in the yield strength, tensile strength and plasticity from 57 to 71 MPa, from 155 to 201 MPa, and from 11.5 to 18.8%, respectively.

The effect of aluminum oxide nanoparticles introduced into the melt using ultrasonic treatment of the melt on the structure, and the properties of pure aluminum were studied in [3]. It was found that the introduction of nanoparticles of aluminum oxide using ultrasonic treatment of the melt into commercially pure aluminum allowed microstructure refining and reduced the average grain size from 200 to 69 μm. The introduction of nanoparticles made it possible to increase the hardness from 19 to 22 HB, the yield strength from 12 to 27 MPa, and the tensile strength from 48 to 79 MPa in commercially pure aluminum.

In [4], the effect of porosity on the properties of aluminum composites reinforced with SiC particles was studied using 2D models and finite elements analysis. The authors found that when the particle content did not exceed 11%, and the theoretical density of the material was reached, the shape of the particles (circular or square) did not affect the mechanical properties of the composite. With SiC particle content of more than 11%, the angular particles were more effective in improving the mechanical properties. Furthermore, it was shown that the particles transferred the stress to the soft matrix even in the presence of pores in the composite.

The analysis of the effect of processing the Al-Mg-Si solution on the transformation of β-AlFeSi particles into α-(FeMn) Si and the aging of the Al 6063 alloy were presented in [5]. It was found that the maximum hardness of 104 HV was reached following a solution treatment at a temperature of 600 ◦C for 2 h and aging at a temperature of 160 ◦C for 12 h.

In work [6], the Al-Si-Cu-Mg alloy was cast at casting speeds of 1, 2, 3 and 4 mm/s by the Ohno continuous casting (OCC) process. It was shown that the spacing between secondary dendrite arms of α-Al dendrites in the samples decreased significantly with an increase in the casting speed. Moreover, an increase in the casting speed to 4 mm/s allowed the authors to increase the tensile strength of the alloy, which was not possible with the alloy obtained at the casting speed of 1 mm/s.

In work [7], the authors obtained the Al8Zn7Ni3Mg hypereutectic alloy, combining an Al-8% Zn-3% Mg matrix reinforced by an Al3Ni intermetallic compound. The specified structure of the alloy was obtained by changing the cooling rate in the range from 0.1 K/s to 2.3 × 105 K/s. It was shown that an increase in the cooling rate made it possible to grind intermetallic compounds in the alloy, the size of which was 50 nm, at a cooling rate of 105 K/s. At the same time, the hardness of the alloy increased to 220 HV.

In work [8], the aluminum alloy containing 0.6 wt.% Zr, 0.4% Fe and 0.4% Si was obtained by electromagnetic casting at a high cooling rate (140 K/s). The subsequent procedure of direct cold drawing was used to synthesize Al3Zr nanoparticles in the alloy with a uniform distribution in the matrix and an average size not exceeding 10 nm. The composite wire sample obtained by the authors had a tensile strength and electrical conductivity of 234 MPa and 55.6 IACS, respectively.

A study of the process of treating commercially pure titanium by friction with stirring using a tool made of a nickel-based superalloy ZhS6U was presented in [9]. It was found that the transfer layer in titanium contained chemical elements of the ZhS6U alloy. After such treatment, the tensile strength of commercially pure titanium increased by 25%.

In work [10], an experiment was carried out on friction stir welding of a sheet made of a fine-grain Al-Mg-Sc-Zr alloy to study the peculiarities of the plasticized metal flow and microstructural evolution. It was shown that the stir zone macrostructure might contain either a single or many nugget zones, depending on sheet thickness and the seam length. Despite the finer-grained structure, the hardness of the welded seam was lower than the hardness of the base alloy, and the tensile strength of the welded seam was comparable to the tensile strength of the base alloy.

The structural phase characteristics of the friction-treated Al-Cu alloy zone were studied in [11]. The presence of Al2Cu, Al2Cu3, AlCu3, Al2MgCu, etc. intermetallic phases with a nonuniform distribution over volume and size was revealed.

In work [12], the structural and phase states of the TiAl system alloyed with rare earth metals were studied. Tantalum, yttrium and dysprosium were used as additives. It was found that the studied systems Ti(49 at.%)–Al(49 at.%) with additions of Ta, Y and Dy contained intermetallic compounds consisting of AlTi3, TiAl in hexagonal, tetragonal and triclinic forms.

### **3. Conclusions and Outlook**

Future development of new casting methods and techniques for local control of crystallization of light alloys will significantly change the approach to industrial production of high-tech and efficient products. This will make it possible to create light alloys for specific structural elements of automobiles, airplanes, etc., where a set of specific operational properties is required. There is also great interest in joining new light alloys, including those with a composite structure, which also indicates the widespread use of new light alloys.

Implementation of this concept will create a strong demand for innovations in the field of new light alloys and their casting and joining technologies, which are one of the leading topics in this special issue. As a guest editor of this special issue, I am very pleased with the final result and hope that these articles will be useful to both researchers and technologists working in the production and processing of light alloys. I would like to express my gratitude to the authors for their contributions and the reviewers who helped with the review process. I would also like to extend a special thanks to the Metals Editor Office staff for working on the special issue.

**Conflicts of Interest:** The author declares no conflict of interest.

### **References**

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### **The Structural and Phase State of the TiAl System Alloyed with Rare-Earth Metals of the Controlled Composition Synthesized by the "Hydride Technology"**

**Akbayan Belgibayeva 1,2,\*, Yuri Abzaev 1,3, Natalia Karakchieva 1, Rakhmetulla Erkasov 2, Victor Sachkov 1 and Irina Kurzina 1**

Received: 30 May 2020; Accepted: 22 June 2020; Published: 29 June 2020

**Abstract:** The structural state and the quantitative phase analysis of the TiAl system, alloyed with rare-earth metals synthesized using hydride technology, were studied in this work. Using the Rietveld method, the content of the major phases in the initial system Ti(50 at.%)–Al(50 at.%), as well as Ti(49 at.%)–Al(49 at.%), with alloying additions Ta, Y and Dy having a high accuracy was determined. The methods of scanning electron microscopy, transmission electron microscope and X-ray spectral microanalysis of the local areas of the structure for studying the distribution of alloying elements were used. The energies of lattices of separate phases were also determined after the full-profile specification. All the lattices of the identified structures (about 30) turned out to be stable. It was established that in the Ti(49 at.%)–Al(49 at.%) systems under study with alloying additions of metals Ta, Y and Dy, there were intermetallides composed of AlTi3, TiAl in the hexagonal, tetragonal and triclinic units. It is known that after microalloying alloys by Y and Dy metals, the mass fraction of TiAl phases increases significantly (>70%).

**Keywords:** intermetallides; phase composition; microstructure; hydrides; TiAl system

### **1. Introduction**

Intermetallic alloys based on γ-TiAl are a good example of the way basic and applied research along with industrial development can lead to obtaining a new innovative class of advanced structural materials [1–4]. Nowadays, intermetallic alloys based on the γ-TiAl phase are promising materials for application in aeronautical engineering, owing to their attractive properties: high specific strength, stiffness, creep resistance at temperatures of *T* = 600–800 ◦C, oxidation resistance and burn resistance at temperatures up to *T* = 900 ◦C. In the temperature range of *T* = 20–800 ◦C, the specific modulus of elasticity of these alloys is higher than that of the nickel by 30–50% [2]. It is supposed that in the gas-turbine engine, light γ-TiAl alloys (ρ ≈ 4 g/cm3) will partially replace heat-resisting heavy nickel alloys (ρ = 8–8.5 g/cm3), which will allow significant increases in its specific power characteristics during a simultaneous decrease in the fuel consumption, carbon dioxide emissions and nosiness [4].

To achieve a certain combination of properties, it is necessary to optimize the chemical composition and the microstructure [4–8]. Therefore, in the past two decades, the increased focus of researchers of γ-TiAl alloys has been on achieving an optimal combination of mechanical properties by varying the elemental composition and the microstructure with various sizes of columns/grains and thickness of plates. For that, a detailed work on the optimization of the composition and conditions of thermomechanical/thermal treatment of the alloys is being performed.

In modern materials science, a new method of synthesis of binary and multicomponent alloys is of interest, which is called a "hydride" technology (HT) [9]. Modern technologies of production of alloys (mechanical alloying, arc smelting, powder metallurgy, etc.,) are associated with notable labor intensity and hardware difficulties (application of a deep vacuum and creation of inert environment at high temperatures, duration and multiplicity of processes, etc.). The powder metallurgy technique is characterized by a special duration because the result of the interaction of metals in the initial mixtures is mainly determined by solid diffusion rates. The specific difficulties of obtaining quality alloys are also associated with the presence of a dense oxide film on the surfaces of particles of refractory metals, which prevents the passage of mutual diffusion. The HT method allows the avoidance of the majority of them (in particular, it excludes melting). HT is a high-tech method based on the combination of self-propagating high-temperature synthesis SHS of transition metal hydrides and heat treatment of a mixture of hydrides, resulting in alloy formation [10–12]. In the HT method, to obtain high-strength alloys based on transition metals, the powders of hydrides of refractory metals and alloys are used as initial materials. The essence of HT consists of the successive use of processes of hydrides synthesis, their combined compacting and dehydration. The advantage of HT is that the alloy is formed at a relatively low temperature (from 900 to 1150 ◦C) during rapid exposition (from 1 to 2 h). It is important that metals with different values of melting temperature and density should alloy without melting [10–12]. Other advantages of the method are relative cheapness, the use of refractory materials, as well as obtaining materials with high purity. It is also known that the formation of metal alloys can be carried out easier from hydrides than from metals themselves, since chemical bonds in hydrides are less strong than bonds in metal structures [11,12].

It is well known that the compaction of the alloy structure is one of the most effective ways of improvement of strength and plasticity of the materials. Alloying was demonstrated as a viable approach to enhance the mechanical properties of TiAl alloys at room temperature through improving the materials microstructure [13,14].

To improve the mechanical properties of γ-TiAl alloys, their alloying with rare-earth elements is of interest. It is known that [15–17] the introduction of rare-earth elements (La, Er and Dy) into titanium and intermetallic alloys based on γ-TiAl can lead to the improvement in machinability in the as-cast state, owing to refinement of the structure, to increased heat-resistance and refractoriness, and in the γ-TiAl alloys, in some cases, to the enhancement of the technological plasticity at elevated temperatures.

Previous studies supposed that the addition of yttrium [18] and gadolinium [19–22] could significantly reduce the grain sizes and lamellar spaces of TiAl alloys, enhance strength and plasticity at room temperature, as well as creep resistance at high temperatures. However, the mechanisms of microstructure refinement in the rare-earth metals-modified (REM-modified) TiAl alloys were studied insufficiently because of the complex succession of solidification and solid-phase transformations of TiAl during thermal treatment and cooling [23]. Systematic data on the influence of additions of rare-earth elements on the structure and mechanical properties of γ-TiAl alloys in the literature are practically not presented.

Thus, the development of the methods for producing high-strength alloys of the TiAl system, the study of the dependence of the structural phase state on the synthesis parameters, and the study of the effect of various modifying additives on the physical and mechanical properties is an actual task.

The purpose of the present work is the development of the basics of hydride technology for the production of TiAl alloys and the study of the influence of alloying additions Ta, Y and Dy on the microstructure and the phase composition of the titanium–aluminum system obtained by the HT method. The Ti(50 at.%)–Al(50 at.%) system has been accepted as a basis of the composite material with REM additions of no more than 2 at.%.

### **2. Materials and Methods**

### *2.1. Obtaining Alloys*

Samples with the following selected atomic compositions were prepared TA (TiAl); TAT (TiAlTa); TAY (TiAlY); TAD (TiAlDy). Sample TA was prepared as follows: a weighed amount of the titanium was placed in a quartz boat and heated in a furnace (Nabertherm RS 120/750/13, Lilienthal, Germany) in a stream of the hydrogen. The heating rate of the furnace was 10 ◦C/min to 450 ◦C, with a hydrogen volume flow of 500 mL/min.

This sample was cured for 3 h at this temperature; after that, it was cooled down to room temperature. The obtained metal hydrides were mixed with a nanodispersed aluminum powder (the average size of particles was 115 ± 10 nm, specific surface area—19.4 ± 3 m2/g, loading of active aluminum—80.8 ± 0.6%). Then, the mixture was pressed into a round plate and was formed (collapsible compression mold, Lab Tools, diameter of 13 mm, thickness of 2 mm), pressing load was under a pressure of 5.3 tons/sm2, bulk density was 3 g/sm3 (Lab Tools PAH-20, 2019, St. Petersburg, Russia)). The equation for obtaining the TA sample can be written as:

$$\text{Ti} + \text{H}\_2 = \text{TiH}\_2 + 27.3 \text{ kcal/mole} \tag{1}$$

The TAT sample was prepared as follows: a weighed amount of the tantalum was placed in a quartz boat and heated in a furnace (Nabertherm RS 120/750/13, Lilienthal, Germany) in a stream of the hydrogen. The heating rate of the furnace was 10 ◦C/min to 550 ◦C, with a hydrogen volume flow of 500 mL/min. This sample was cured for 3 h at this temperature; after that, it was cooled down to room temperature. The equation for the production of titanium hydride can be written as:

$$\text{2Ta} + \text{H}\_2 = \text{2TaH} + \text{56 kcal/mole} \tag{2}$$

Titanium hydride was obtained as in the case of sample TA. Then, the obtained titanium hydride and tantalum hydride were mixed with a nanodispersed aluminum powder (according to the manufacturer, the average size of the particles was 115 ± 10 nm, specific surface area—19.4 ± 3 m2/g, loading of active aluminum—80.8 ± 0.6%). Then, the mixture is pressed similar to TA.

Samples of TAY and TAD were obtained similarly to the method of obtaining TAT. The temperature for yttrium hydride and dysprosium hydride was 420 ◦C. The equation for obtaining a TAY sample can be written as:

$$\text{Y} + \text{H}\_2 = \text{Y} \\ \text{H}\_2 + 56 \text{ kcal/mole} \tag{3}$$

$$\text{2Y} + \text{3H}\_2 = \text{2YH}\_3 + \text{21.5 kcal/mole} \tag{4}$$

The equation for the preparation of dysprosium hydride can be written as:

$$2\text{Dy} + 3\text{H}\_2 = 2\text{DyH}\_3\tag{5}$$

The obtained sample blanks were obtained in a vacuum unit and heated to a temperature of 1150 ◦C, with a heating rate of 2 ◦C/min. The vacuum value was 10−4 atm. They were maintained under these conditions for 3 h and cooled at a rate of 5 ◦C/min.

The equations for obtaining samples TA, TAT, TAY and TAD can be written as:

$$\text{TiH}\_2 + \text{Al} = \text{TiAl} + 2\text{H}\_2\uparrow\tag{6}$$

$$2\text{TiH}\_2 + 2\text{Al} + 2\text{TaH} = 2\text{TiAlTa} + 3\text{H}\_2\uparrow\tag{7}$$

$$\text{TiH}\_2 + \text{Al} + \text{YH}\_2 = \text{TiAlY} + 2\text{H}\_2\uparrow\tag{8}$$

$$2\text{TiH}\_2 + 2\text{Al} + 2\text{DyH}\_3 = 2\text{TiAlDy} + 5\text{H}\_2\uparrow\tag{9}$$

### *2.2. Research Methods*

The structural state and the quantitative phase analysis of the system TiAl-alloying metal (TA-REM), synthesized by HT, were studied in this work by the Rietveld method and scanning electron microscopy (SEM) [24,25]. Ta, Y and Dy metals were used as alloying additions. The X-ray diffraction studies of the TA-REM system were undertaken using DRON4-07 (Bourevestnik, St. Petersburg, Russia), which was modified for the digital processing of the signal. Spectra were made using copper radiation (Ka) according to the Bragg–Brentano scheme with an increment of 0.02◦, exposure time at a point of 1 sec and in the angle range of 10–86◦. The voltage on the X-ray tube was 30 kV, beam current—25 mA. The structural state and the quantitative content of the phases were identified by the Rietveld method by means of reflex [26–29]. As the standard lattices, the crystallographic data of the COD base [26] and the model structures of the TiAl system, predicted by the program code USPEX with the interface shell SIESTA [30], were used. In connection with the promising physicochemical properties of the TiAl systems, they were alloyed with metal additions [24–34]. In this work, the search for standards of the TiAl system was supplemented by the USPEX program with the outer shell SIESTA [31,35,36].

To study the distribution of alloying elements in the structure at the local level, the quantitative X-ray spectral microanalysis of the near-surface layer, accompanied by the analysis of the microstructure and the morphology of the surface by the method of scanning electron microscopy, was carried out using the X-ray fluorescence spectral analyzer S4 Pioneer (Bruker, Karlsruhe, Germany) and the raster electron microscope "LEOEVO 50" (Zeiss, Oberkochen, Germany).

Electron microscope studies of the surface microstructure of TA-REM samples were conducted using the transmission electron microscope "JEM-2100F" (JEOL, Tokyo, Japan) with accelerating voltage of 200 kV using the attachment "JEOL" (JEOL, Tokyo, Japan), intended for energy-dispersive spectral analysis. The phases were identified applying well-known procedures using schemes of microdiffraction patterns calculated by the table values of parameters of crystalline lattices.

### **3. Results and Discussion**

The USPEX-SIESTA program by means of the evolution code can predict stable structures of the known elemental composition, space group, among which the structures with the global minimum of enthalpy of the system are predicted. In the present work, the TiAl structures were used for the quantitative phase analysis. For the discovered structures (about 30) based on complete structural information, the energies of lattices were calculated proceeding from the first principles. All the lattices turned out to be stable and they were used at the next stage for qualitative and quantitative phase analysis by the Rietveld method. In Tables 1–5, there are standards in the initial state, obtained from the COD base (AlTi3, AlTi), as well as the structure predicted by means of the USPEX code (TiAl-Struct2, AlTi3-2768). For the alloy with the conventional designation TiAl-Struct2, AlTi3-2768, the energy was calculated in the framework of the electron density functional by the gradient pseudo potential of electron density (GGA). The details of the code are given in the work of reference [26]. The total energy of the lattices was determined at 0 K. Wave functions of valence electrons of TiAl-Struct2 atoms were analyzed on the basis of flat waves with a cutoff radius of kinetic energy of 330 eV. In this case, the convergence of the total energy was ~ 0.5 × 10−6 eV/atom. It was established that the lattice energy after geometric optimization turned out to be equal to ETiAl-Struct11 = −4978.555 eV, and in the initial state −4976.099 eV. The final enthalpy of the TiAl-Struct2 structure in both cases is evidence of its stability. TiAl lattices (AlTi3-2768) with Y, Ta and Dy additives (Tables 2–4) also turn out to be stable, and the lattice energies are substantially negative. In this work, based on the results of the qualitative phase analysis of the TA-REM system, it was suggested that the REM nanoadditions were embedded into the interstitial site [0.5 0.5 0.5] of the lattice of the AlTi3 alloy. Quantum chemical calculations of the AlTi3 energy in the initial state, as well as with the embedded nanoadditions in the mentioned site, were made. It was established that: EAlTi3 = −19,318.285 eV; EAlTi3-Y = −19,712.792 eV; EAlTi3-Ta = −19,613.620 eV; EAlTi3-Dy = −31,227.561 eV.

 **%**

**Table 5.** Relative coordinates of the atoms in the TiAl-Struct2 lattice.

The calculations are evidence of the fact that the introduction of the alloying metals into the indicated interstitial site is possible and it leads to the significant stabilizing effect of the lattices in each TA-REM system without exception. In fact, the binding energy of (|EAlTi3|) atoms in the lattices accompanied by the addition of the alloying metal increased significantly. It is interesting to note that the increase in the binding energy is accompanied by significant polarization of the Millikan charges [37,38]. The Millikan charges in the metal additions are equal: [(−5.96)Y], [(−4.42)Ta], [(+0.73)Dy], but on the atoms of the main elements—[(−0.16)Al, (+1.540)Ti], [(−0.35)Al, (+1.22)Ti], [(−0.13)Al, (−0.12)Ti], [(−0.17)Al, (−0.13)Ti] in ternary compounds TAT, TAY and TAD, respectively. The analysis of distances between atoms of the main elements and additions showed that these distances were significantly less than the sum of the covalent radii of free elements Ti and Al, and Ta, Y and Dy, which are equal, respectively, to 1.6, 1.21, 1.7, 1.90, 1.92Å [39]. For instance, in the TAT system, the lengths of bonds TiTa and AlTa are equal to 2.030 and 2.031 Å, respectively. Polarization of the charges is indicative of the growth in the share of the covalent bond as a result of introduction of the alloying metals.

The results of the quantitative content of phases in the studied systems are given in Tables 1–4. The tables show the phases, the structural state of the phases before (Init) and after (Spec) the Rietveld refinement of the experimental diffractograms, the corresponding lattice parameters, the lattice volume (V), the space group, the fraction of the individual phases (Share) and the lattice energy after refinement by the Rietveld method (E). The tables also show the fraction of the calculated integrated intensity (Rwp) in the experimental diffractogram. The analysis of the contributions into the integral intensity of separate phases (Table 1, Figure 1) showed that to a high degree of reliability (Rwp < 7.2%), the major phases were intermetallides AlTi3, AlTi3-REM, AlTi and TA-Struct2 in the initial and optimized states. The experimental diffraction patterns of the TA, TAT, TAY and TAD systems can be approximated well by the calculated integral intensity; the difference between them is an insignificant value (Figure 1). However, as Tables 1–4 show, the contributions of separate phases differ in different systems. Using the Rietveld method, it is established that in the TA system, the contribution of the triclinic TiAl system (over 73%) predominates. Both lattices differ by parameters, atoms coordinates, as well as the volume (Table 1). The number of the main systems also includes the hexagonal lattice of AlTi3 (22.33%). The introduction of an insignificant amount of Ta, Y and Dy leads to the redistribution of the contributions of the hexagonal and triclinic systems. In the TA-REM systems, the contribution of the geometrically optimized lattice (TA-Struct2-GeomOpt) disappears, and the phase with the AlTi tetragonal lattice appears (Tables 2–4). The introduction of Ta into the AlTi3 lattice significantly increases its binding energy and stimulates the contribution to the integral intensity up to 26.99%. On the contrary, in the TAY and TAD systems, additions Y and Dy lead to the reduction in the share of AlTi3 to 16.66, 4.78 and 11.20%, respectively. The share of phases and the energy of the lattices allow evaluation of the effective energy of the TA, TAT, TAY and TAD systems according to the formula αE1+βE2+γE3, where α, β, γ—the share of phases; E1, E2, E3—the energy of the lattices of separate phases. It was established that the energies of the mentioned systems were equal to −7907.453, −5407.31, −7341.104, −21,188.219 and −15,920.023 eV, respectively.

**Figure 1.** 3D lattice of the AlTi3 alloy with the Y atom embedded in the site [0.5 0.5 0.5].

The relative coordinates of the atoms and isotropic parameters of displacements are given in Tables 5 and 6). The lattice parameters and space group are shown in Tables 1–4.

**Table 6.** Relative atomic coordinates in the AlTi lattice (AlTi3-2768) with an embedded atom Dy, Ta and Y.

Figure 2 shows the results of the observed and calculated intensities of the initial alloy and alloys obtained by doping with rare-earth metals. X-ray phase analysis of the samples obtained by doping with rare-earth metals showed that they have a complex multiphase structure: the initial components (Ti, Al) and the new phases—intermetallic phases AlTi3, AlTi3–REM, AlTi, TiAl-Struct2—in the initial and optimized states were identified.

**Figure 2.** Diffraction patterns of alloys TA (**a**,**b**); TAY (**c**,**d**); TAT (**e**,**f**); TAD (**g**,**h**): 1—experiment, 2—summary model intensity, 3—difference between intensities, 4–6—contributions to the integral intensity of separate phases.

The study of the microstructure showed that the base of the material was the alloy of titanium and aluminum with a small quantity of alloying metals (no more than 2 at.%). The ratio between the atomic concentrations of Al and Ti is close to the initial 1:1 (Figure 3). The spectra of characteristic X-ray radiation are also given in Figure 3.

As a result of the quantitative analysis, it was established that the matrix of the TAY alloy on average had a composition in mass percent: 29.60% of Al and 68.35% of Ti, which corresponds to the intermetallide phase α2-Ti3Al according to the stoichiometric ratio (Figure 3a). The elemental analysis of the TAT and TAD alloys (Figure 3b,c) also revealed the presence of titanium aluminide α2-Ti3Al. This agrees with the results obtained in the works [40,41]. It was established that the microstructure of the rest of the alloys was also characterized by the formation of the α2-Ti3Al intermetallide.

The choice of these metals as alloying microadditions in the alloys under study is conditioned by a number of positive effects. In particular, metals as an effective modifier of the cast structure [15–22] possess an increased affinity with oxygen, which in turn leads to a substantial decrease in the amount of oxygen in the alloy and, as a consequence, to an increase in low-temperature plasticity because of the reduction in the number of barriers in the form of oxygen atoms, which decelerate the dislocations motion during deformation.

The X-ray phase analysis results of the initial intermetallic alloy TA showed that the basic thermodynamically stable phases were intermetallide compounds Ti3Al, TiAl, TiAl2 and a solid aluminum solution in titanium α-Ti (Table 7).

**Figure 3.** EDX-spectra and elemental composition of the TAY (**a**), TAT (**b**) and TAD (**c**) alloys.

**Table 7.** Crystallographic data of phases in the system TA.

A small fraction of the sample consists of the following intermediate phases having practically constant chemical compositions: Ti5Al11, Ti3Al5, TiAl2, Ti2Al5. The X-ray phase analysis (XPA) results are confirmed by TEM studies. The TEM studies showed that, before the introduction of rare-earth metals, in the initial state, the TA alloy contained phases Ti3Al, TiAl, TiAl2, α-Ti, Ti5Al11 and Ti3Al5 (Figure 4e). Consequently, the obtained TA alloy has a complex multiphase structure, as it contains a number of phases with different crystal lattices. This fact confirms a possibility of obtaining intermetallic alloys by hydride technology.

**Figure 4.** Electron microscope images of the matrix of the initial alloy TA: (**a**) bright-field image; (**b**) microdiffraction pattern; (**c**), (**d**) dark-field images in reflexes, marked by arrows; (**e**) microdiffraction pattern identification.

First of all, the obtained studies of the microstructure confirmed the formation of Ti3Al and TiAl phases. The Ti3Al phase is an ordered phase with a D019 superstructure having an HCP crystal lattice; its space group is *P63*/*mmc*. Ti3Al is formed as parallel lamellar precipitates (Figure 4a). The intermetallide Ti3Al and TiAl phases under formation are microdimensional, which is evident owing to the dark-field images (Figure 4c,d). The presence of Ti3Al and TiAl is associated with

thermodynamics of the phases' formations. The formation of these phases is characterized by a minimum of Gibbs energy [42]. The existence of these intermetallide phases must result in significant alloy strengthening [6,43,44].

A layer-by-layer location of intermetallide phases—that is, alternation by the composition—is possible. Thus, using hydride technology allows for obtaining of complex sandwich structures. Owing to alternation of intermetallide phases, it becomes possible to obtain super strong materials, where each layer strengthens the previous one.

To predict the composition of the formed phases when microalloying the TiAl alloys with rare-earth metals, ternary diagrams of the corresponding ternary systems were considered. The analysis of the TiAlY ternary diagram of the partial isothermal cross-section at 1000 ◦C shows that the content of the Ti:Al:Y = 49:49:2 components corresponds to the highlighted area of formation of γ-TiAl + YAl2 phases (Figure 5). The ternary diagram of the TiAlTa system at 1100 ◦C implies that at a ratio of Ti:Al:Ta = 49:49:2 components the TiAl phase is formed (Figure 5c). According to the ternary diagram, TiAlDy at a ratio of components Ti:Al:Dy = 49:49:2, the formation of the area, where DyAl2 + TiAl phases are present, can be expected (Figure 5b).

**Figure 5.** Isothermal cross-sections of the TiAlY system at 1000 ◦C [45] (**a**), of the TiAlDy system at 500 ◦C [46] (**b**), of the TiAlTa system at 1100 ◦C [47] (**c**).

The X-ray phase analysis of the samples, obtained during alloying with yttrium, is given in Table 8.

**Table 8.** Crystallographic data of phases in the TAY system.

According to the XPA data, the TAY alloy contains the basic Ti3Al5 phase. The Ti3Al5 phase has a rhombic crystal lattice and the space group P/mmm. Along with the Ti3Al5 grains, the structure of the alloy contains a small number of grains of the Ti3Al phase, having an HCP crystal lattice and the space group P63/mmc. A small fraction belongs to phases TiAl, Al, α-Ti and Y.

Figure 6 presents electron microscope images of the TAY alloy. On the microdiffraction pattern (Figure 6b), there are reflexes belonging to planes (110) and (201) TiAl; (101), (200), (302), (411) Ti3Al; (400) Ti3Al5; (311) YAl2; (100) α−Ti; (200) Al. Thus, the XPA results are confirmed by TEM.

**Figure 6.** Electron microscope images of the TAY alloy: (**a**) bright-field image; (**b**) microdiffraction pattern; (**c**,**d**) dark-field images in reflexes, marked by arrows; (**e)** microdiffraction pattern identification.

Figure 7 presents TEM images of the depositions of yttrium particles of different shapes inside the grain of titanium aluminide. According to the TEM results, yttrium and Al form the YAl2 intermetallide in intermetallide phases and are located on dislocations or on grain boundaries.

**Figure 7.** TAY alloy microstructure.

The X-ray phase analysis of the samples obtained when alloying with dysprosium (TAD) and tantalum (TAT) showed that these alloys had a complex multiphase structure: initial components Al, α-Ti and new phases—intermetallide phases TiAl, Ti3Al, Ti5Al11, Ti3Al5—were identified (Table 9). The basic phase of the TAD alloy is the TiAl phase, which has an ordered tetragonal-distorted face-centered structure and the P4/mmm space group.

**Table 9.** Crystallographic data of phases in the TAD and TAT systems.

The results of the X-ray phase analysis of the TAT alloy showed that in the alloy, the main share belonged to intermetallide phases: 30.1% of titanium aluminide TiAl of tetragonal syngony with the space group P4/mmm; 22.9% of the Ti3Al phase, having a space-centered lattice of tetragonal syngony with the space group P63/mmc; 26.4% of the Ti3Al5 phase with the space group P/mmm (Table 9). The smallest share belongs to 1.9% of Ti2Al5 with the space group P4/mmm and 3.9% of Ti5Al11 with the space group P/mmm. In addition, the intermetallide of tantalum and aluminum (Ta161,8Al282,2) is present in the sample.

The results of the X-ray phase analysis of the TiAlDy alloy are confirmed by the TEM results (Figure 8). On the microdiffraction pattern there are reflexes belonging to planes (002), (004), (402), (413) TiAl; (600), (203) and (423) Ti3Al; (312) TiAl2; (008) Ti5Al11; (511) Al; (300) α-Ti; (551) DyAl2.

Figure 9 shows the electron microscope images of the TAT alloy. According to the microdiffraction analysis in the layers of the TAT alloy, there are the following intermetallide phases with reflexes belonging to planes (312), (412) TiA; (110), (402) and (210) Ti3Al; (311), (440) Ti3Al5; (314) Ti5Al11; (103) α-Ti; (008) and (215) TaAl3.

**Figure 8.** Electron microscope images of the TAD alloy: (**a**) bright-field image; (**b**) microdiffraction pattern; (**c**) dark-field images in the reflex; (**d**) microdiffraction pattern identification.

**Figure 9.** Electron microscope images of the TAT alloy: (**a**) bright-field image; (**b**) microdiffraction pattern; (**c**) dark-field image in reflex; (**d**) microdiffraction pattern identification.

Thus, the TEM studies confirmed the formation of a number of phases of titanium aluminide. The tantalum compounds indicated by the XPA method probably remain in the volume of the sample, but only TaAl3 is present in the film structure.

As Ti and Ta have significant solubility in Al, Ta can form a solid substitutional solution in titanium and contribute to formation of the Widmanstatten microstructure. In addition, it is well known that the TaAl system also dissolves Ti (Ti aluminide) by as much as 60 at.%. Aluminides TiAl and Ti3Al dissolve Ta (Ta aluminide) up to 10 and 15 at.%, respectively [47–49]. The TEM research results comply with the data of the X-ray phase analysis, which imply that an insignificant amount of Ta dissolves in the crystals of TiAl and Ti3Al and are present in the (Ti,Ta)Al3 phase.

Thus, the XPA and TEM studies have shown that the formed solid solution based on titanium and the alloys under study have particles of the new phase in their structure. It is also well known that after microalloying of the alloys with metals Y and Dy, the mass fraction of the TiAl phases increases significantly (>70%). A distribution of tantalum in the titanium–aluminum matrix is evidence of the possibility of the formation of the three component system (Ti,Ta)Al3 [47–49]. Alloying metals, Y and Dy, form a compound with aluminum; they are applied as a dispersion phase and strengthen the alloy structure by the introduction of the second phase into the metal matrix in the form of a high-melting compound. If YAl2 is located on the dislocations or on the grain boundaries, DyAl2 is chaotically distributed in the grain volume.

### **4. Conclusions**

The composite materials based on gamma-aluminides of titanium TiAl-REM were obtained by a new method—"hydride technology". The content of the major phases in the initial TA system, as well as TA-REM, containing microadditions Ta, Y and Dy, was determined by the Rietveld method with high reliability.

1. It was established that in the initial TA system and in the TA-REM system, containing microadditions, TiAl predominate in tetragonal and triclinic units. It was established that the used initial standard lattices, as well as the structures after the full-profile specification, were in the stable states.

2. The results of X-ray phase analysis are confirmed by the results of TEM. It is known that after microalloying alloys by Y and Dy metals, the mass fraction of TiAl phases increases significantly (>70%).

3. From the results of X-ray phase analysis and TEM studies, it follows that an insignificant amount of Ta dissolves in the crystals of TiAl and Ti3Al and are present in the (Ti,Ta)Al3 phase. This fact indicates the possibility of the formation of a three component system (Ti, Ta)Al3.

**Author Contributions:** Y.A. studied structural state and determined the energies of lattices of separate phases and fraction of quantitative phase, and wrote the paper; N.K. conducted the microstructure research of samples and analyzed the data; I.K. conducted the X-ray phase analyses of samples; A.B. wrote the introduction and conducted the synthesis of samples; V.S. and R.E. analyzed data. All authors have read and agreed to the published version of the manuscript.

**Funding:** This work was carried out with the support of the Program of Competitive Recovery of TSU, NU <IP_ADDRESS>8 L. and project No. 10.3031.2017/4.6.

**Acknowledgments:** The authors are grateful to Mark Kalashnikov for fruitful cooperation.

**Conflicts of Interest:** The authors declare no conflict of interest. The funding sponsors had no role in the design of the study; in the collection, analyses, or interpretation of data; in the writing of the manuscript, and in the decision to publish the results.

### **References**

© 2020 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

### *Article* **Microstructure of In-Situ Friction Stir Processed Al-Cu Transition Zone**

### **Anna Zykova, Andrey Chumaevskii, Anastasia Gusarova, Tatiana Kalashnikova, Sergei Fortuna, Nickolai Savchenko, Evgeny Kolubaev and Sergei Tarasov \***

Institute of Strength Physics and Materials Science, Siberian Branch of Russian Academy of Sciences, Tomsk 634055, Russian<EMAIL_ADDRESS>(A.Z<EMAIL_ADDRESS>(A.C<EMAIL_ADDRESS>(A.G<EMAIL_ADDRESS>(T.K<EMAIL_ADDRESS>(S.F<EMAIL_ADDRESS>(N.S<EMAIL_ADDRESS>(E.K.) **\*** Correspondence<EMAIL_ADDRESS>Tel.: +7-3822-286-863

Received: 22 May 2020; Accepted: 16 June 2020; Published: 18 June 2020

**Abstract:** The majority of literature sources dedicated to dissimilar Al-Cu friction stir welding testifies to the formation of intermetallic compounds (IMC) according to diffusion-controlled reactions, i.e., without liquation on the Al/Cu interfaces. Fewer sources report on revealing Al-Cu eutectics, i.e., that IMCs are formed with the presence of the liquid phase. This work is an attempt to fill the gap in the results and find out the reasons behind such a difference. Structural-phase characteristics of an in-situ friction stir processed (FSP) Al-Cu zone were studied. The single-pass FSPed stir zone (SZ) was characterized by the presence of IMCs such as Al2Cu, Al2Cu3, AlCu3, Al2MgCu, whose distribution in the SZ was extremely inhomogeneous. The advancing side SZ contained large IMC particles as well as Al(Mg,Cu) solid solution (SS) dendrites and Al-Al2Cu eutectics. The retreating side SZ was composed of Al-Cu solid solution layered structures and smaller IMCs. Such a difference may be explained by different levels of heat input with respect to the SZ sides as well as by using lap FSP instead of the butt one.

**Keywords:** in-situ friction stir process; aluminum alloys; Al-Cu metallomatrix composite; intermetallic compounds; diffusion-controlled reactions; Al-Cu eutectics

### **1. Introduction**

Metallomatrix composite surface materials modified using friction stir processing (FSP) are state-of-the-art materials, which are intended to combine high strength, wear resistance with high ductility and fatigue resistance of a core metal such as an aluminum alloy [1]. The FSP was originally a process used for surface structural modification, i.e., grain refining, which then was adapted to prepare the metallic matrix composites (MMC) surface coatings by means of introducing various reinforcement particles and admixing them to the matrix metal [1–10]. The FSP utilizes the friction-generated heat for plasticizing the matrix metal, which then is transferred to the rear zone by means of tool rotation and translational motion. The plasticized metal flows along a rather complex trajectory and its adhesion to the tool plays a great role in the metal transfer and stirring. The intense stirring serves to ensure homogeneous distribution of hard particles throughout the stir zone (SZ), and the FSP process parameters such as tool rotation rate and travel speed can be varied to find an optimal degree of mixing as well as temperature conditions.

The hard particles may be introduced into the stirred metal directly [7–9,11–13] or form in-situ inside the metal [14–18] by means of solid-state reactions between the admixed components, between the admixed components and the matrix or between the dissimilar metals processed. The FSP preparation of hybrid composites with the use of in-situ reactions shows up some advantages over those obtained using commercially available reinforcement particles. The first advantage is that in-situ reactions allow us to obtain finer and more homogeneously distributed reinforcing particles [17]. The second advantage is that either coherent or semicoherent boundaries may form between these in-situ prepared particles and the matrix [10,15,19] and therefore more thermodynamically stable and strong particle/matrix interfaces are feasible [20]. The literature sources were analyzed to show that many experimental compositions such as Al7075-Ti-6Al-4V [14]; Al1050-Ni-Ti-C [15]; Al-SiC [15,21]; Al6061-Fly ash [18]; Al1050-Fe2O3-Al [17]; Al-1050-Cu [16]; Al-Ni; Al-Nb [22]; Al-graphene [13,21,22] were subjected to FSP in order to prepare the in-situ MMCs.

Commonly, copper is one of the widely used and efficient metals to carry out FSP on aluminum alloys and thus obtain the MMCs reinforced with Al-Cu intermetallic particles [23]. Intermetallic Cu-Al particle reinforced MMCs may be fabricated using friction stir welding (FSW) on dissimilar metals [24–27], multilayer friction stir brazing [28], die-casting, powder metallurgy, etc.

The effect of different tool configurations on friction heat generation, metal flow and formation of intermetallic layers was studied in friction stir spot welding of dissimilar Cu and Al metals [24]. It was shown that thin intermetallic layers were formed from compounds such as CuAl2, CuAl and Al4Cu9. Intermetallic compound layers consisting of CuAl2 and Al4Cu9 were found at the Cu-Al boundaries along with hot cracking during butt friction stir welding of Cu and Al [25].

AA6061-T6 plates were welded with a copper plate inserted in the butt line between two AA6061 plates and the resulting stir zone structure contained CuAl2 and Al4Cu9 intermetallic particles [26]. Both intermetallic particles and interlayers were formed during butt FSW on Cu and Al plates [27]. Ultrasound treatment was used to reduce the thickness of intermetallic layers formed with FSW [29].

Only a few publications were devoted to obtaining the Al-Cu composites using the FSP admixing the Cu powder. Hsu et al. [30] demonstrated that homogeneous Al-Al2Cu MMC with Young modulus 88 ± 8 GPa, yield stress 450 MPa, ultimate stress 650 MPa and 0.15 plasticity may be obtained using FSP on Al-15 at.% Cu green samples compacted at 225 MPa.

It has been reported [16] that when pure copper powders were FSP admixed to aluminum plates at the tool travel speed 1.66 mm/s and rotation rate 750 rpm, the resulting intermetallic particles were CuAl2 ones. Increasing the FSP pass number and, in particular, the FSP tool rotation rate to 1000 and even to 1500 rpm resulted in precipitation of mainly Al-Cu and Al4Cu9 intermetallics. On the contrary, only Al2Cu precipitates were found in the aluminum irrespective of the FSP pass number [31].

This work was focused on studying the specificity of in-situ synthesis of Al-Cu intermetallic particles by means of lap FSP on a copper and aluminum alloy.

### **2. Materials and Methods**

The hot-rolled AA5056 sheets were cut into 200 mm × 60 mm × 5 mm samples. C11000 copper sheets were cut into 200 mm × 60 mm × 2 mm samples. Chemical compositions of the alloy and copper sheets are shown in Table 1.

**Table 1.** Chemical composition of A5056 and C11000 plates.

Friction stir processing was carried out with the use of the FSW machine (Sespel, Cheboksary, Russian) at the Institute of Strength Physics and Materials Science of Siberian Branch of Russian Academy of Sciences (Tomsk, Russian) (Figure 1). A truncated cone flute FSW tool with a 2.5 mm height pin and top and bottom diameters of 6 and 4 mm, respectively, was used. The tool shoulder diameter was 12 mm. The FSW tool inclination angle was 3◦. The FSW parameters were as follows: rotation rate 500 rpm, travel speed 90 mm/min, plunging force 12,000 N. This set of parameters was found to be optimal as follows from previous experimenting [32]. The FSW tool penetration was 2.5 mm.

**Figure 1.** Scheme of FSP on C11000/A5056 sandwich (**a**) and single-pass FSP seams (**b**).

The FSPed samples were cut using electric dischage method (EDM) in planes perpendicular to the joint centerline to obtain specimens for examination and tests (Figure 1a). The microstructural evolution was examined on polished and etched section views prepared according to ASTM standards and visualized using optical microscopes Altami Met 1S (LLC Altami, Sankt Petersburg, Russian) and Olympus LEXT 4100 (Olympus NDT, Inc., Waltham, MA, USA) as well as scanning electron microscopy (SEM) and transmission electron microscopy (TEM) instruments Zeiss LEO EVO 50 (Carl Zeiss, Oberkochen, Germany) and JEOL-2100 (JEOL Ltd., Akishima, Japan), respectively. The chemical composition of precipitates was controlled using an EDS attachment to TEM.

The mean particle sizes were determined using the linear intercept method. Perfect stoichiometric compound component ratios were used to identify intermetallics found in the stir zone and analyzed with EDS. Using an X-ray diffraction (XRD) instrument XRD-7000S (Shimadzu, Kyoto, Japan) operated at 35 kV, 24 mA, irradiation was applied for identifying the Al-Cu phases. Microhardness profiles were obtained using a microhardness tester Duramin 5 (Struers A/S, Ballerup, Denmark) at 100 g load and a dwell time of 10 s.

### **3. Results**

The macrostructure view of the FSW seam cross-sectional area allows us to observe composite structures in both the stirring zone (SZ) and thermomechanically affected zone (TMAZ) (Figure 2a). Both central and bottom areas of the SZ located close to the advancing side of the seam reveal discontinuities which may be shrinkage pores formed during the formation of Cu-Al intermetallics. The stir zone is characterized by alternating Al/Cu layers in its bottom part (Figure 2b) as well as Al-Cu solid solution and intermetallic compound (IMC) layers (Figure 2c). The large IMCs areas are seen on the retreating side of the SZ (Figure 2a).

The XRD pattern in Figure 3 reveals the phases as follows: Al, Cu, Al2Cu, Al2Cu3, AlCu3, Al2MgCu and thus suggests that in-situ Al-Cu reactions have occurred. Nevertheless, there are large IMC-free areas composed of unreacted Cu and Al. Formation of Al2Cu IMCs on the Cu/Al interfaces during FSW was noted in the majority of works dedicated to dissimilar FSW [24–27,29,30].

**Figure 2.** The FSPed SZ cross-sectional zones and Cu-Al intermixed areas: 1—defects; 2—solid solution (SS) areas; 3—IMC areas; 4—cracking at the advancing side; AS—advancing side; RS—retreating side; (**a**) the SZ macrostructure; (**b**) alternating SS and IMC layers; (**c**)—aluminum inflow pattern.

**Figure 3.** The XRD pattern of the Cu-Al stirring zone metal.

The composite structure of Al-Cu SZ does not look structurally homogeneous since it includes many different phases and microstructures (Figure 4).

The IMC layers formed at the Al/Cu interfaces on the Cu-rich side of SZ are composed of 2 to 5 μm in size Al2Cu3 and AlCu3 IMCs (Figure 4a,d). The mean size of the Al2Cu3 and AlCu3 particles in these IMCs is about 300 nm. The presence of these phases is confirmed by the results of XRD (Figure 3), TEM (Figure 5) and EDS (Table 2, spectra 4–7).

**Figure 4.** The SEM BSE images of SZ microstructures after a single-pass FSP on the Al-Cu sandwich: (**a**) SS/IMC interface; (**b**) stir zone macrostructures; (**c**) column IMCs; (**d**) IMCs of different compositions; (**e**) substrate/stir zone boundary; (**f**) mixed IMC + SS zone.

**Figure 5.** TEM images of SZ areas: Al-Cu interfaces with eutectics (**a**,**b**); EDS profile across the Al2Cu/Cu transition zone (**c**); SAED pattern obtained from the area in Figure 5a and SAED reflection identification (**d**); 1–15 are the EDS probe spots corresponding phases shown in Table 2.

**Table 2.** Chemical compositions of phases shown in Figure 5.

The interfaces between fine-crystalline copper and IMCs are shown in Figure 5a as well as Al/Al2Cu eutectics located in between the Al2Cu particles. The selected area electron diffraction (SAED) pattern obtained from the area in Figure 5a shows the reflections, which can be identified as those belonging to AlCu3, Al2Cu and Al2Cu3 (Figure 5d). Analyzing the EDS spectra and taking into account the ideal stoichiometric formulas of the IMCs detected, the theoretical compositions of them were determined and are presented in Tables 2 and 3. It is worthwhile noting that IMCs composed of Al2Cu3 and AlCu3 particles are inherent to all Al/Cu alternating layers (Figure 4a,d–f).

**Table 3.** Chemical compositions of phases shown in Figure 6.

**Figure 6.** The SEM BSE images of microstructures in the SZ (**a**,**b**), which contain SS Al(Cu,Mg), Al2Cu and Al2MgCu particles as detected using the EDS probe on microstructure components denoted 1, 2, 3, 4.

When looking at the Al-rich part of the SZ, more aluminum-rich phases are formed there according to reactions as follows:

$$\text{Al}\_2\text{Cu}\_3 + 5\text{Al} \to \text{Al} + 5\text{Al}\_2\text{Cu} \tag{1}$$

$$\text{AlCu}\_3 + \text{6Al} \rightarrow \text{Al} + \text{3Al}\_2\text{Cu}.\tag{2}$$

The microstructure of this zone is composed of SS Al(Cu,Mg) and Al2Cu particles as confirmed by the EDS and SEM (Figures 4 and 6, Table 3).

The solid solution in this zone contains up to 5.7 at.% Cu and 3.1 at.% Mg in Al, i.e., it can be referred to as an Al(Cu,Mg) phase as shown by the EDS spectra in Figure 6 and Table 3. These Al(Cu,Mg) particles are dendrites and have a mean size of 11.8 μm (Figure 6). Such a dendritic shapes may be evidence in favor of heterogeneous nucleation and growth in a liquid phase.

Two different sorts of CuAl2 particles are formed in the SZ (Figure 4a,c–f; Figure 6) such as thin 7 × 50 μm2 area platelets and the fine ones found in eutectics. Figure 5c shows an EDS profile along the line shown in Figure 5a, i.e., a transition zone from IMC to the Al-Al2MgCu eutectics and then to the recrystallized fine-grained copper.

An intermetallic Al-Al2MgCu eutectic phase was also EDS detected in the middle of SZ (Tables 2 and 3) (Figures 3, 4a and 6).

The macroscopic FSP track cross-sectional area in Figure 7a shows lines along which the microhardness number profiles (Figure 7b,c) were obtained. Both Figure 7b and 7c demonstrate that microhardness profiles obtained in two perpendicular directions allow for the differentiation between the matrix and IMCs, and, in fact, reveal the SZ structural inhomogeneity, which means a lack of equality in strength. The microhardness of IMCs is by a factor of 2 to 5 higher than those of the base metals.

**Figure 7.** Microhardness profile measurement lines (**a**) and microhardness profiles along horizontal (1, 2, 3) (**b**) and vertical (4, 5, 6) directions (**c**).

#### **4. Discussion**

The results of this work clearly show the inhomogeneous structure of the stir zone composed of copper, aluminum, Al-Mg-Cu solid solution and a set of different morphology IMCs. The FSP is a strong nonequilibrium process so that the microstructural evolution of the processed metal is determined by a variety of external factors such as heat generation, mechanical stirring (deformation), heat removal as well as internal process factors such as adhesion-assisted or quasi-viscous transfer of metal portions to the zone behind the tool, dynamic recrystallization, diffusion-controlled precipitation and mechanochemical solid-state reactions. The Cu-Al system is capable of forming intermetallic compounds with a high exothermic effect so that a thin liquid phase layer may form at the Al/Cu interface due to contact melting. Such a phenomenon leads to a fast liquid-phase synthesis of coarse IMC layers and particles, especially when fusion methods are used to obtain the Cu-Al alloys. At the same time, only diffusion-controlled formation of IMCs is possible when the process temperatures are low enough. In FSP on Cu-Al, the temperatures in the stir zone are in the range 400–500 ◦C [31], i.e., lower than the eutectic temperature *T*E = 548.2 ◦C and no Al-Cu eutectics were detected in this work. All IMCs were formed by means of diffusion-controlled precipitation from a supersaturated solid solution obtained in FSP according to the model suggested by Pretorius et al. [33]. According to such a model, the effective heat of formation (Δ*H* ) of phases at the binary Al-Me system interfaces can be determined as follows:

$$
\Delta H\_i' = \Delta H\_i^0 \cdot \frac{\mathbb{C}\_\varepsilon}{\mathbb{C}\_\varepsilon} \tag{3}
$$

where Δ*H i* is the effective heat of formation of *i*-phase, Δ*H*0 *i* is the enthalpy of formation change for phase *i*; *Ce* is the effective concentration of the limiting element at the interface; *Cc* is the concentration of the limiting element in the compound. Taking into account that Δ*G*0 ≈ Δ*H*0 the effective Gibbs free energy change (Δ*G i* ) in case of an *i*-phase formation is determined as:

$$
\Delta G\_i' = \Delta G\_i^0 \cdot \frac{\mathbb{C}\_{\mathfrak{c}}}{\mathbb{C}\_{\mathfrak{c}}}.\tag{4}
$$

Table 4 shows the results of calculating the Gibbs free energy changes corresponding to the formation of all Al-Cu binary system phases. It can be noted that only negative Δ*G i* values were obtained, thus, determining the feasibility of the IMC nucleation and growth. The maximum absolute Δ*G i* values Δ*G* Al2Cu3 = −31.28 kJ/mol and Δ*G* AlCu3 = −22.84 kJ/mol were found for Al2Cu3 and AlCu3, i.e., these phases were the first ones to form at the Cu/Al interfaces. Therefore, those phases were detected in this work as small particles on the Cu-rich side of the image in Figure 4a. The next phase to form was Al2Cu with Δ*G* Al2Cu = −19.54 kJ/mol.

It should be noted that, in general, the diffusion-controlled Al-Cu interaction may lead to the formation of a variety of IMC phases as dependent on Cu-content, sort of source materials (sheet or powders), etc. Since copper and aluminum sheets were used in this work, the local Cu-content may be as high as 40 vol.% due to intense stirring and transfer then three binary phases were formed such as Al2Cu3, AlCu3, and Al2Cu. However, it was reported [16,30,31,34] that only a single Al2Cu phase was formed when FSP admixing the pure copper powders at a concentration of ≤15 at.% and in-situ synthesizing the Al-Cu composite. Using the Pretorius model [33], it was experimentally established [31] that Al2Cu had to be the first phase to form since its Gibbs free energy change of formation was the most negative value of all other phases as determined for a Cu at.% concentration corresponding to the lowest temperature eutectic.

**Table 4.** Effective Gibbs free energy changes at Cu-Al interfaces at 450 ◦C.

On the other hand, eutectic structures were observed during friction stir spot welding of AA5083 to copper [36]. The majority of papers devoted to studying the dissimilar Al-Cu butt FSWed SZ did not show any presence of eutectics in distinction to the friction stir spot welding (FSSW) SZ microstructures. The reason is the higher heat removal into a copper sheet in the case of butt FSW as compared to FSSW or lap FSW. Preferential localization of IMCs on the advancing side of the SZ may be related to more intense admixing between aluminum and copper layers (particles) detached off the parent metals and transferred to the stagnation (trailing) zone behind the tool closer to the advancing side where these layers adhere to the already deposited layers; finally, recrystallize and grow the IMCs at rest. Closer to the retreating side these layers are stirred by the tool and deformed so that the strain dissolution of the IMCs prevails over their precipitation, thus, facilitating the formation of the solid solution only. Such a consideration is based on the adhesion-assisted transfer of metal during FSP when a transferred portion (layer) of metal first adheres to the FSW tool, then is transferred to the stagnation zone behind the tool, and adheres back to the previously transferred layers [37].

It is suggested also that more heat is generated on the advancing side as compared to that of the retreating side [38]. This factor may provide higher heating and better conditions for the contact melting between copper and aluminum transferred layers.

The combined effect of severe plastic deformation and heating provided the formation of a row of Al-Cu phases. There are a number of phases whose morphology allows for the suggestion of their origin from a liquid phase. First of all, those phases were Al-Al2Cu eutectics and Al-Mg-Cu solid solution dendrites that formed near the IMC large particles as a result of depleting these zones of copper. Such a scenario is inherent to the advanced side of the SZ zone.

Reaction-diffusion controlled nucleation and growth of IMC smaller particles is inherent to the mechanically alloyed multilayer metal on the retreating side of the SZ. Friction stir processing in this zone resulted in the formation of a Cu-Al mechanical alloy (bronze) which differed also from both parent metals by its color. The smaller IMCs are found in the interlayer spaces whereas the in the main phase there is the Al-Mg-Cu SS.

### **5. Conclusions**

Microstructural evolution and phase composition of stir zone in-situ obtained using friction stir processing on an Al-Cu bimetal workpiece were studied:

4. The IMC areas containing the eutectics and solid solution dendrites, which might originate from Al-Cu liquation, are characterized by large irregular shaped shrinkage pores.

**Author Contributions:** A.Z., A.C., A.G., T.K., S.F., N.S. performed sample preparation and characterization; A.C., E.K. and S.T. performed project administration, conceptualization, supervision, designed the experiments, and analyzed the data; A.Z., A.C. and S.T. wrote this paper. All authors have read and agreed to the published version of the manuscript.

**Funding:** Sample preparation and processing were funded by the Government research assignment for ISPMS SB RAS, project No. III.23.2.4. Characterization of microstructures and phases obtained using dissimilar FSP on copper and aluminum alloy was performed with financial support from RSF Grant No. 19-79-00136.

**Conflicts of Interest:** The authors declare no conflict of interest.

### **References**

© 2020 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

### **Microstructural Analysis of Friction Stir Butt Welded Al-Mg-Sc-Zr Alloy Heavy Gauge Sheets**

### **Tatiana Kalashnikova \*, Andrey Chumaevskii, Kirill Kalashnikov, Sergei Fortuna, Evgeny Kolubaev and Sergei Tarasov**

Institute of Strength Physics and Materials Science, Siberian Branch of Russian Academy of Sciences, 634055 Tomsk, Russia<EMAIL_ADDRESS>(A.C<EMAIL_ADDRESS>(K.K<EMAIL_ADDRESS>(S.F<EMAIL_ADDRESS>(E.K<EMAIL_ADDRESS>(S.T.)

**\*** Correspondence<EMAIL_ADDRESS>Tel.: +7-3822-286-863

Received: 31 May 2020; Accepted: 15 June 2020; Published: 17 June 2020

**Abstract:** Friction stir welding (FSW) on a heavy gauge sheet of a hereditary fine-grained Al-Mg-Sc-Zr alloy was carried out to study the specifics of plasticized metal flow and microstructural evolution in different sections and zones of the joint. It was found that the stir zone (SZ) macrostructure may contain either a single or many nugget zones depending on the metal sheet thickness and the seam length. The effect of grain kinking in a thermomechanically affected zone (TMAZ) under pressure from the stir zone metal was discovered. The stir zone metal was fine-grained but had a microhardness lower than that of the base metal, which may be explained by the overaging effect of FSW on the Al3Sc precipitates. The tensile strength of the joint was almost equal to that of the base metal (BM). The grain size distributions were obtained in different sections below the sheet surface and away from the exit hole, which allowed us to suggest the specific adhesion-assisted layer-by layer metal transfer mechanism in FSW.

**Keywords:** friction stir welding; aluminum alloy; structure formation; adhesion; metal transfer; mechanical properties

### **1. Introduction**

Al-Mg alloys are well-known structural materials that are widely used in various applications including those where welding is a single option joining method. Improving the Al-Mg alloy strength characteristics is achieved by alloying the metals with scandium and zirconium [1–4], which makes them hereditary fine-grained alloys while retaining their corrosion resistance [2,5].

The alloy fine-grained microstructure is stabilized by the precipitation of nanosized Al6Mn, Al3Sc, and Al3(Sc,Zr) matrix coherent particles, which first precipitate from a solid solution and then have a pinning effect on the grain boundaries so that even annealing at 723 K for 16 h does not increase the grain size [6]. The Al3Sc precipitates are the most effective precipitates for pinning the grain boundaries [7]. It is known also [8] that the nanosized Al3Sc precipitates are stable against coalescence so that that their mean diameter increased only from 13.5 to 15.4 nm for 7 days at 350 ◦C.

Friction stir welding (FSW) is a method used for joining metals in a solid state and is, therefore, often used on metals and alloys that cannot be welded using any fusion welding technique. An Al-Mg-Sc-Zr alloy is an example of such a hardly weldable alloy that and is, therefore, a candidate for joining with friction stir welding. It is common knowledge that this solid state joining method usually forms a fine-grained microstructure in the stirring zone. The presence of the nanosized growth inhibitors in Al-Mg-Sc-Zr is an extra factor that helps retaining the fine-grained structure of the welded metal and thus improving its mechanical characteristics [9]. For example, it was reported [10] that FSW increased the ultimate and yield stresses of the Sc + Zr-alloyed aluminum alloy by 23.8 and 11.9%, respectively. Practically a 100% tensile strength was achieved on the Al-5.4Mg-0.2Sc-0.1Zr alloy FSW seam even

after annealing [11]. The mechanical strength advantage of the FSW joint compared to that of the TIG on the Al-Mg-Sc-Zr alloy was thus demonstrated [10].

Another positive factor of the FSW-generated fine-grained structure is its potential to take advantage of superplasticity characteristics [12–16] that allow up to 2000% elongation while retaining microstructural thermal stability at 450 ◦C [12].

The effect of Al3Sc and Al3(Sc,Zr) nanosized precipitates on pinning the fatigue crack tip was also reported [17]. It was shown [18] that friction stir processing improves the corrosion resistance of Al-Mg alloys in a NaCl aqueous solution.

An important problem in FSW is obtaining a flawless and reliable seam when joining thick sheets. The inhomogeneity and efficiency of metal stirring, enhanced adhesion of metal to the FSW tool, diffusion-controlled FSW tool wear, fast heat removal, and high mechanical loads are factors that aggravate the problems commonly encountered with FSW on aluminum alloy sheets up to 10 mm thickness. Some technologies require welding heavy gauge sheets or pipes, which then are then machined to obtain the final structure. This requirement makes it necessary to study FSW seams on heavy gauge workpieces.

It may be concluded, therefore, that Al-Mg-Sc and Al-Mg-Sc-Zr alloys not only possess high mechanical characteristics and are weldable using FSW but also that their characteristics can be further improved using FSW, which is a promising research field from both practical and scientific points of view.

The severe thermomechanical impact of FSW on the plasticized aluminum alloy in the stirring zone creates the conditions for the strain-induced dissolution of particles, thereby enriching the solid solution with the alloying elements and facilitating the fast precipitation of particles. This is an element of FSW aspect that has not received much attention in the literature, especially when taking into account the dynamic character of dissolution/precipitation, as well as recrystallization, during the process. This issue becomes even more intriguing when applying FSW to an Al-Mg-Sc-Zr alloy, whose microstructure is known to have high temperature stability. The FSW joint structure formation is determined by the specificity of metal transfer and the role played by the adhesion of plasticized metal to the FSW tool.

The objective of this paper is to study the microstructures and mechanical characteristics of the FSW joint zones obtained on heavy gauge Al-Mg-Sc-Zr alloy sheets.

### **2. Materials and Methods**

Hot-rolled 30 and 35 mm thick 300 mm × 300 mm sheets of AA1570 and HSS M2 steel FSW tools were used to carry out friction stir welding using an FSP machine PowerStir 345C (Holroyd Precision Ltd., Milnrow, UK) at the facilities of S.P. Korolev RSC Energia Experimental Machine-building Plant ZEM. The shape and parameters of the tools used for the welding are shown in Figure 1. The process parameters were adjusted experimentally to obtain the FSW joints, as shown in Figure 1. The FSW tool rotation rate, traverse speed, and plunging force were 600 RPM, 10 mm/min, and 40 kN, respectively. The FSW seam can be divided into three zones as shown in Figure 1a, where zones B, S, and H denote the FSW tool plunging zone characterized by insufficient heating plasticization, steady welding, and the FSW tool exit zone, respectively.

The stop-action technique [19] was applied to study the mechanical behavior of the metal in the thermomechanically affected zones (TMAZs) when the FSW tool exit hole was sectioned and then marked as shown in Figure 2. Another series of metallographic views was obtained by EDM cutting the seam by planes normal to the sheet surface and both perpendicular and parallel to the FSW joint line. An EDM cutting machine DK7745 (Suzhou Simos CNC Technology Co., Ltd. Suzhou, China) was used to cut off the samples, which were then ground, polished, and etched to prepare the metallographic section views.

**Figure 1.** The friction stir welding (FSW) seam zones on a 35 mm thick AA1570 sheet (**a**), the FSW tool used for the 35 mm thick sheet after welding (**b**), the FSW tool used for the 30 mm thick sheet after welding (**c**), and the FSW tool used for the 35 mm thick sheet before welding (**d**). 1–tensile specimens; 2–metallographic section views, dmax–the larger pin diameter, h–the pin height, α–the taper angle.

The joint strength was tested using a test machine BISS UT-04-0100 (BISS (P) Ltd. Bangalore, India) on samples cut off the seam by a plane perpendicular to the tool travel direction at different distances below the top surface so that the stir zone was located in the middle of the sample's gauge length part (Figure 1). The loading speed was 100 mm/s.

A microstructural examination was performed using optical microscopes Altami MET 1T(Altami Ltd., Saint-Petersburg, Russia), Olympus OLK41102 (Olympus Corp, Tokio, Japan), SEM and TEM instruments Tescan VEGA 3 (TESCAN ORSAY HOLDING, BRNO, Czech Republic), Zeiss LEO EVO 50 (Carl Zeiss AG, Oberkochen Germany), and JEOL JEM-2100 (JEOL Ltd., Akishima, Japan).

Thin foils for TEM were prepared using EDM cutting, mechanical polishing, and ion thinning to represent the stir zone metal in two perpendicular sections with respect to the joint line. A microhardness tester Duramin-5 (Struers A/S, Ballerup, Danemark) was used to obtain microhardness profiles across the FSW seam zones at a 490 mN load.

**Figure 2.** A diagram showing the sample cut-off scheme: (**a**) the leading edge of the exit hole on the retreating side (LE-RS), (**b**) the leading edge of the exit hole on the advancing side (LE-AS), (**c**) the trailing edge of the exit hole on the retreating side (TE-RS), (**d**) the trailing edge of the exit hole on the advancing side (TE-AS).

### **3. Results**

### *3.1. Stir Zone and Thermomechanically A*ff*ected Zone Macrostructures*

The transverse cross-sectional view in Figure 3 reveals the inhomogeneous macrostructure of the deep penetration FSW joint stir zone. The onset part of the joint (zone "B" in Figure 1) is characterized by a dense small nugget zone formed in the root part of the seam (Figure 3a,b), while "loose" metal with discontinuities (Figure 3a,b, defect 1) is formed in the top part. The FSW area behind the tool is characterized by the presence of a crack (Figure 3a,b, defect 2).

**Figure 3.** The macrostructures of the FSW joint on 35 mm thickness alloy sheet in section parallel to the joint line: (**a**) the cross section of zone B (**b**); the cross section of zone S (**c**); the multi-nugget macrostructure of the stir zone (**d**); and a cross section view of the joint on the 30 mm thickness metal (**e**).

The steady welding zone "S" is characterized by another type of macrostructure, where numerous nugget structures are formed in the stir zone (SZ) instead of a single nugget (Figure 3c,d). These nugget structures are evidence of change in the metal flow compared to that of zone "B". The top part of the stir zone obtained on the 35 mm thick sheet still contains pores and discontinuities.

This type of macrostructure is also typical for the FSW joint obtained on the 30 mm thick sheet (Figure 3e), with the exception of a wormhole defect formed in the half-height AS zone because this 30 mm thick sheet was made of two thinner ones.

The FSW tool exit zone H revealed the metal flow patterns formed by the tool around the exit hole. Figure 4a shows the microstructure of sample 4.2 cut off the trailing part of the exit hole, as shown in Figure 2. The bottom metal shows a large and dense nugget (Figure 4a, SZ1), above which a number of metal flow patterns (SZ2–SZ7) can be seen. The "loosest" macrostructure is observed on the SZ8 pattern, which is formed by a shoulder-driven metal flow.

**Figure 4.** Macrostructures in the FSW tool exit zone as viewed from the plane parallel to the welding direction and perpendicular to the sheet's surface: (**a**) trailing part of the exit hole, sample 4.2; (**b**) the exit hole on the advancing side of the weld under different magnifications (**c**,**e**); kinked TMAZ grains in the top (**f**) and bottom (**d**) parts, respectively.

The macrostructure of the AS exit hole quadrant planar section (Figure 4b) (cut off as shown in in Figure 2, section 4.1) is represented by a difficult-to-see stir zone and a TMAZ that can be divided into the ultrafine grained TMAZ I in the vicinity of the SZ and TMAZ II, which features larger kinked grains located closer to the base metal (BM) (Figure 4c,d). Only a small portion of the SZ can be observed using this view. The rolling welding defect between the two sheets is also clearly seen here (Figure 4b). The TMAZ zone structures show orientation dependence with respect to the different exit hole quadrants. This is also true for such structural characteristics as grain size and grain shapes (Figure 4c,e).

Since the base metal hot-rolled grains were preferentially oriented with their longest axes along the rolling direction and their boundaries were pinned by the nanosized precipitates against continuous recrystallization, mechanical pressure in TMAZ II exerted from the SZ distorted their shapes. Distinct from the TMAZ I superfine grains, whose deformation at the SZ/TMAZ boundary was fully determined by the friction force, the TMAZ II grains oriented with their long axes along the pressure retained their sizes but instead experienced kinking. The alternating kinking band patterns can be delineated as shown in Figure 4d and look similar to those generated with the shear banding mechanism. Those grains whose shorter axes coincide with the pressure direction were more resistant to kinking. Kinking is also more pronounced in the bottom parts of the TMAZ II where the FSW tool-induced metal flow pressure is higher.

The planar section views in Figures 5–7 represent the structures formed at different distances below the metal surface. The structures formed close to the leading edge of the exit hole in sections 2.2–2.6, as well as those formed at the trailing edge of the exit hole in sections 3.2–3.6, are shown in Figures 5 and 6, respectively.

**Figure 5.** The microstructural evolution on stir zone (SZ), thermomechanically-affected zone (TMAZ), heat-affected zone (HAZ) and base metal (BM)microstructures near the exit hole leading edge on the advancing side (AS), as seen in plane sections 2.2 (**a**), 2.3 (**b**), 2.4 (**c**), 2.5 (**d**), and 2.6 (**e**) (Figure 2).

**Figure 6.** The alloy structures in the exit hole leading edge on the AS side, as viewed in sections 3.2 (**a**), 3.3 (**b**), 3.4 (**c**), 3.5 (**d**) and 3.6 (**e**) (Figure 2).

The leading AS edge planar sections 2.2 to 2.5 (see Figure 2) are represented by a very narrow stir zone (Figure 5a–d) but a wide TMAZ composed of both elongated and almost equiaxed grains, whose origin may be explained by horizontal plane sectioning of the kinked grains. The TMAZs are wide in sections 2.4–2.4 (Figure 5a–c), while those in section 2.5 (Figure 5d) demonstrate a narrow TMAZ that becomes wider again in section 2.6 (Figure 5e).

The grain size in the TMAZ is gradually increased from the SZ/TMAZ to TMAZ/HAZ boundary (Figure 8).

When examining the structures and zones formed on the trailing edge of the hole in sections 3.2–3.6 (Figure 6), it can be noted that the TMAZ is wider in the RS trailing edge (TE-RS) (Figure 6) of the exit hole than the TMAZ in the leading edge AS part (LE-RS) (Figure 5).

**Figure 7.** The TMAZ grain structures in planar sections 3.2 (**a**), 3.3 (**b**), 3.4 (**c**), 3.5 (**d**) and 3.6 (**e**), as viewed normal to the sheet plane.

These differences are clearly seen in Figure 7, where the TMAZ grain structures are shown at a higher magnification. sections 3.2 to 3.4 reveal moderately elongated grains whose length is about twice as much their width. sections 3.4 and 3.6, respectively, demonstrate almost equiaxed and extremely elongated thin grains.

**Figure 8.** The TMAZ grain size distributions as measured along lines 3 (**a**), 2 (**b**), and 1 (**c**) (in Figures 5 and 6 (**a**,**c**)) and the dependence of the TMAZ size on the height of the 30 mm thick sample (**d**). The diagram labels are in accordance with those of the diagram in Figure 3.

#### *3.2. Stir Zone and Thermomechanically A*ff*ected Zone Microstructures*

The specifics of the TMAZ grain structure modification in the exit hole area are shown in Figure 8. Transversely to the SZ RS, along line 3 (shown in Figure 6), the change of the grain size from the stir zone to the heat-affected zone at different depths is determined by two main characteristic features. In samples from the upper part of the welded joint zone (sections 3.2, 3,3, and 3.4 in Figure 2), there is an initial rapid growth of grains with a subsequent drop and a further growth of the grain size, which becomes equal to the grain sizes of the heat-affected zone (HAZ) or the base metal. In samples from the bottom part of the joint (sections 3.5 and 3.6), the grains gradually grow from the stir zone to the heat-affected zone (Figure 8a). There is no grain growth near the stir zone. This feature indicates different deformation and recrystallization processes at different distances from the surface of the FSW-joint. At the leading edge of the tool exit hole on AS along line 2 (shown in Figure 5), the SZ-boundary material in sections 2.2, 2.3, and 2.4 is characterized by larger grains compared to those in sections 2.5 and 2.6 (Figure 8b). At a distance of more than 5 mm from the stir zone, the TMAZ grain size is leveled on all samples.

Along welding line 1 (Figure 5), the increase in the grain size from the stir zone to the heat-affected zone occurs gradually on all samples (Figure 8c). Significant differences are observed only for sections 2.5 and 2.4, far from the stir zone.

Unevenness in the size of the TMAZs can be observed both in different directions (for the welding axis) and at different distances below the weld surface (Figure 8d). It is expected that the smallest size of the TMAZ will be characteristic of the leading edge, where the TMAZ size changes smoothly and identically along line 1 and line 2. In the trailing edge area on the retreating side and across the stir zone along line 3, the size of the stir zone increases from the minimum in the bottom of the joint to the average in the central volume, with the maximum found in the shoulder-affected area. These changes show strain inhomogeneity in different areas of the joint.

The microstructures of the SZ zone in both joints are represented by recrystallized grains and coarse Al3Sc precipitates (Figure 9a,b). It can be seen from Figure 9c,d that the coarse incoherent particles cannot effectively pin the grain boundaries, so they continue migrating through the precipitates leaving chaotically distributed dislocations and dislocation loops.

**Figure 9.** TEM images of the stir zone coarse precipitates, grains, and dislocation loops in 35 mm (**a**,**c**) and 30 mm (**b**,**d**) welds.

TEM images of the TMAZ microstructures in both FSW seams are characterized by the presence of numerous extinction contours inside the elongated and bent grains (Figure 10a, shown with arrows). The presence of several bend extinction contours inside an elongated grain is evidence of alternating sign flexural strain (corrugation). These corrugated grains are geometrically subdivided into shorter subgrains to accommodate the bending. Some of these subgrains are separated by high-angle boundaries from the neighboring subgrains (Figure 10b, shown with arrows).

**Figure 10.** TEM images of the TMAZ structures in the 35 mm (**a**) and 30 mm (**b**) thickness seams.

### *3.3. Mechanical Strength*

The results of the mechanical tests show that the welded metal has a strength close to that of the base metal (Figure 11a). The minimal tensile strength of 235–280 MPa was obtained when testing the samples cut off the top and bottom sections 1.1, 1.2, 1.9, and 1.10 (see Figure 3e).

Tests in the defect-free zone "S" (Figure 1) areas show that the ultimate tensile strength (UTS) of the specimens is about 350–390 MPa (Figure 11b). The exception is Specimen 2.10, cut from near the facing surface of the sample, whose UTS is about 305 MPa. The dependence of tensile strength on the distance from the weld root is shown in Figure 11c.

Testing of the welded samples shows their high defectiveness due to the formation of large discontinuity in the weld, which is detrimental to the metal's mechanical properties (Figure 11c,d). Samples from the defect-free areas have strength properties at the level of equivalent samples from 35 mm thick welded plates at a level of 400–405 MPa. In the presence of defects, this strength is very low (even close to zero). The difference between zone "S" and zone "B" (Figure 1) is seen in the diagrams. Only one sample cut from the center zone showed a strength close to zero, while about one third of them cut from zone "B" had similar results.

When comparing the mechanical properties in the cross sections of materials of both thicknesses, the defect localized in the structure becomes clearly distinguished (Figure 11e,f). For joints made of 35 mm-thick plates, the poor strength shown on samples from the upper part is typical (Figure 11e). For samples made of 30 mm-thick plates, the strength loss is typical for samples from the central part of the joint (Figure 11f). It should be noted that, generally, samples demonstrate rather low mechanical properties in comparison with that of the base metal, which could be caused by over-ageing of the SZ material during welding. This process is characterized by accelerated precipitation of the particles from the solid solution, making the particles lose their coherence with the matrix and thus impairing the joint mechanical characteristics. At the same time, the mechanical properties of the base metal also have a significant heterogeneity when testing samples cut off along and across the rolling direction (Figure 11g). In this case, the samples cut along the rolling direction have a slightly higher strength and less ductility than those cut across the rolling direction.

**Figure 11.** Mechanical properties of specimens cut across the weld at different distances from the weld root for 35 mm-thick (**a**,**b**) and 30 mm-thick (**c**,**d**) plates, the dependence of the mechanical properties on the distance from the weld root for 35 mm-thick (**e**) and 30 mm-thick (**f**) plates in zones "B" and "S" (see Figure 1), and the test diagrams of the base metal tested along and across the rolling direction (**g**).

The discontinuous yielding effect, strongly pronounced during the testing of the base metal specimens (Figure 11g), is eliminated during the testing of the FSW joined samples. During the deformation of specimens along and across the welding axis, the most pronounced manifestation of the Portevin Le Chatelier effect is observed in cases where the tensile axis coincided with the welding axis. This leads to high amplitude serrations under the tensile loading. In samples of the weld material, the phenomenon of discontinuous yielding was observed practically on all specimens cut from the

defect-free zones as well as on a small part of the defective ones (Figure 11a–d). Some specimens demonstrated deformation behavior similar to that of the base metal ones tested along the rolling direction, and others showed the same with specimens tested across the rolling direction. For example, curves a1, a2, and a3, shown in Figure 11d, demonstrate similar changes in strain to those in diagrams of the specimens cut along the rolling direction rd1–rd3, while curves 2.1 and 2.3 are similar to those in diagrams of the specimens cut along the transverse direction td1–td3. The main difference between the above groups of specimens is their differently pronounced stress changes during deformation. In this case, the strength of the base metal in the tensile test along the rolling direction is 405 MPa, which coincides with the strength of the 30 mm-thick joint specimens in the transverse direction of 365 MPa, which is close to the strength of the defect-free specimens of the 35 mm-thick joint.

The microhardness profiles obtained from both joints demonstrate the reduced microhardness numbers of the stir zone (Figure 12) compared to those of base metal. Such a finding is typical for heat treatable aluminum alloys because this zone is usually over-aged in FSW such that the precipitates lose their coherence, and dispersion hardening becomes ineffective. No extra hardening was noted in the heat-affected zone, which may be explained by the thermal stability of both grains and precipitates in this zone.

**Figure 12.** The microhardness profile across the FSW joint at 5 mm below the surface on 35 mm thickness metal.

### *3.4. Metal Transfer*

The macrographs in Figures 13 and 14 show grain refining in the base metal via friction from the rotating FSW tool. It was suggested earlier [20] that metal transfer in FSW may be performed in a layer-by-layer manner when the grain-refined and plasticized metal layer first sticks to the FSW tool surface and is then transferred to the zone behind the tool and deposited there on a previously deposited layer. The stop-action technique allows one to observe plasticized metal layers, which are separated from the bulk of the metal on the advancing side of the hole (Figure 13), transferred to the retreating side (Figure 14), and finally deposited on the previously transferred and bonded layers.

**Figure 13.** Optical images of the microstructures formed in the FSW tool exit hole metal leading edge metal (section 2.2 in Figure 2). Blue dotted line contour denotes the transfer layer area.

**Figure 14.** Optical images of the microstructures formed in the FSW tool exit hole metal trailing edge metal (section 3.2 in Figure 2). The blue dotted line contour denotes the transfer layer area.

SEM BSE images allow one to distinguish by contrast the intermetallic precipitates of solid solution grains (Figure 15). The majority of these coarse precipitates look like streaks oriented along the rolling direction (as well as tool travel speed) and are clearly seen in the base metal, while fewer are distributed chaotically.

**Figure 15.** SEM BSE composite image of the metal and precipitates in the FSW tool exit hole leading edge metal (section 2.2 in Figure 2).

The friction force induced deformation not only distorts the grains but also changes their orientation, which can be observed by the orientation of the precipitate streaks (Figure 15).

Completely recrystallized metal is observed in a thin 0.5–0.6 mm layer, which was stuck to the FSW tool and thereby transferred in a spiral bottom-to-surface trajectory.

### **4. Discussion**

The stir zone formation in high thickness materials may occur in a very complex multiflow manner so that a number of vortex structures with different configurations forms there, partially due to the low pressure in the welding zone. As a result, a welded joint has, in addition to structural irregularities, significant differences in its mechanical properties with respect to the distance below the top surface. Despite the presence of structural irregularities, it is possible to accurately characterize the process of the welded joint formation in the tool pin stir zone (excluding the area of shoulders influence). As described above, the direct material transfer by the tool is predefined by the primary subdivision of the base metal grains through severe plastic deformation with the formation of a wide thermomechanically affected zone (Figure 5). The grain size is expected to increase with distance from the stir zone to the heat-affected zone (Figure 8). The different widths of the TMAZ at different levels of the weld indicates that the process of grain subdivision and, consequently, the formation of the material transfer layer (Figures 14 and 15) occurs with a different intensity depending on the distance below the top surface. This ensures that the stir zone consists of several nuggets and makes the material transfer frequency potentially differ from level to level. This irregularity in the adhesive

interaction also likely affects the mechanical characteristics. However, the peculiarities of the formation of structures at each of the studied weld levels demonstrate that the welded joint formation in the tool pin stir zone occurs via the same mechanism.

The results obtained in this work, as well as those obtained earlier, suggest that FSW seam formation occurs according to the scheme shown in Figure 16.

**Figure 16.** Schematic diagram to illustrate the adhesion-assisted transfer of metal with the FSW pin. 1—tool, 2—rotation direction, 3—primary grain subdivision, 4—formation of transfer layer, 5—reverse transfer (deposition), 6—successive layer-by-layer metal transfer and deposition.

After plunging the FSW tool into a metal, the tool starts travelling along the joint line. The plasticized and grain refined metal layer is transferred by the rotating tool to the trailing zone where it is detached from the tool and sticks back to the hot metal deposited. This cycle is repeated again and again to form the welded seam. The transferred layer thickness is determined by the temperature and shear stress/strain gradients exerted from the tool into the base metal. This model was inferred from numerous experiments on unlubricated and adhesion controlled sliding friction when direct and reversed metal transfer occurs between the contacting bodies [20–22].

This discontinuous transfer mechanism is responsible for the formation of inhomogeneous grain structures, SZ nuggets, and the TMAZ. An overly strong adhesion of the plasticized metal to the FSW tool may have detrimental effects on the stirring and metal flow efficiency. The transferred portion of metal may not detach from the tool surface and thus be involved in the next cycle, thus breaking the regularity of transfer and forming additional metal flows, which can then provide the final multi-nugget pattern shown in this work in Figure 3c,d. This is especially likely in FSW on thick metal sheets with high pressure and tool travel resistance.

On the other hand, overly weak adhesion would not allow the metal to form a welded joint at all. Therefore, the adhesion of the metal to the tool is an important factor in metal transfer, which has never been considered in modeling. It is known, for example, that it is possible to obtain an FSW seam using an absolutely smooth FSW pin surface [23,24] and that it is the adhesion that works here to help transfer the metal.

Our results show that the SZ metal precipitates are not coherent with the Al-Mg matrix; thus, grain structure recrystallization is not stopped and continues by means of grain boundary migration and diffusion-controlled dissolution of the incoherent precipitates, leading to numerous dislocation loops forming in the metal. It is known that the dislocation loops can appear by a vacancy condensation mechanism in diffusion on the grain boundaries. It seems that the coarse precipitates almost have no pinning effect of the grain boundaries so that grains may grow as large as several micrometers. Severe deformation and heating during FSW facilitates strain dissolution of the initial precipitates that, consequently, have to precipitate again. Intense stirring in the SZ dissolves them again and again until stirring is stopped and the metal comes to rest. This is where precipitation overtakes dissolution, and precipitates grow very quickly on the defective structure. That is why the coarsest precipitates are found in the TMAZ (Figure 3), where the strain is low and some particles are hiding there and thus avoid the strain dissolution. In addition, Figure 8a shows that some grain growth can occur in the vicinity of the SZ/TMAZ boundary. This finding can be explained by the same reason, i.e., that the coarse incoherent precipitates are formed in the TMAZ and allow the grains to grow. Nevertheless, even if the precipitates are semicoherent, they may still retain some pinning effects on the grain boundaries.

The TMAZ zone is characterized by kinked grains (Figure 4d), which are the results of deformation under the pressure exerted by the FSW tool. A similar effect was observed under compression deformation of the Al-Mg-Sc alloy (AA5024) at 523 K [25]. TEM mages show that these kinked grains are composed of even smaller subgrains with high-angle boundaries (Figure 13) that are plausibly formed by microshears in the band-like grains, according to the scheme proposed by Belyakov et al. [26]. Another possible mechanism could be geometrical recrystallization [27], which is considered to be a type of continuous recrystallization. It was reported [28] that continuous recrystallization is inherent with the Al-Mg-Sc-Zr alloys.

#### **5. Conclusions**

As a result of the study, it was found that the stir zone of FSW joints on heavy gauge Al-Mg-Sc-Zr alloy sheets is a complexly structured zone with the presence of several basic structural elements. At the beginning of the welded joint, there is a nugget with an onion ring structure of a small size, and a highly defective area is located closer to the weld top surface. In the middle and end of the 35 mm-thick joint, there is a stir zone with a complex structure that consists of a large number of vortex or turbulent material flows formed by its movement along the tool contour.

In samples of both thicknesses, defects can occur throughout the entire weld length and reduce the strength of the joint area to almost zero. The macro-defect formation localization is possible both in the lower third of the thickness in the advancing side of the weld and in the upper part in the stir zone in the shoulder effect area. Alternatively, it can be located on the border of the stir zone in the shoulder-driven area and in the pin-driven area. In the zone of defect formation, the strength properties are about 235–280 MPa for 35 mm-thick plates and for 30 mm-thick plates in the defect areas, the strength decreases until reaching zero, which might be caused by different types of formed defects. The mechanical properties of the defect-free material are thus at a sufficiently high level. Specimens from the defect-free areas of the 30 mm-thick joints have strength properties around 400–405 MPa, while specimens from 35 mm-thick welded joints have a lower strength of about 350–390 MPa. The strength of the base metal during the tensile test along the rolling direction is 405 MPa, and in the transverse direction it is 365 MPa. In terms of strength and discontinuous yielding effect implementation, specimens from the 30 mm-thick plates were closer to the base metal tested along the rolling direction, while specimens from the 35 mm-thick plates were closer to the material tested across the rolling direction.

The structure of welded joints is formed by the successive layer-by-layer transfer of a material by means of adhesive interaction. Due to the high thickness of the plates to be welded, the process of joint formation becomes more complicated, and the layer transferred by the tool is divided into separate sections with the presence of an upward motion component along with the motion around the tool. The high thickness of the Al-Mg-Sc-Zr alloy plates being welded might affect the plastic deformation

process during the FSW so that metal flow was rearranged into separate small vortex zones, along with the formation of complex material flows along the tool contour and, as a result of the above, the formation of a multi-nugget stir zone with a complex structure.

**Author Contributions:** Conceptualization, writing—original draft preparation, methodology, investigation, visualization, T.K. and A.C.; formal analysis, writing—review and editing K.K.; investigation, S.F.; project administration, supervision, E.K.; conceptualization, S.T. All authors have read and agreed to the published version of the manuscript.

**Funding:** The work was performed according to the government research assignment for ISPMS SB RAS, project No. III.23.2.4.

**Conflicts of Interest:** The authors declare no conflict of interest.

### **References**

© 2020 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

*Article*

### **Wear of ZhS6U Nickel Superalloy Tool in Friction Stir Processing on Commercially Pure Titanium**

**Alihan Amirov \*, Alexander Eliseev \*, Evgeny Kolubaev \*, Andrey Filippov \* and Valery Rubtsov \***

Institute of Strength Physics and Materials Science of Siberian Branch Russian Academy of Sciences 2/4, pr. Akademicheskii, 634055 Tomsk, Russia

**\*** Correspondence<EMAIL_ADDRESS>(A.A<EMAIL_ADDRESS>(A.E<EMAIL_ADDRESS>(E.K<EMAIL_ADDRESS>(A.F<EMAIL_ADDRESS>(V.R.)

Received: 29 May 2020; Accepted: 15 June 2020; Published: 16 June 2020

**Abstract:** The use of electric arc or gas welding in the manufacture of titanium components often results in low quality welded joints due to large residual stresses and strains. A successful solution to this problem can be found in the application of friction stir welding. However, friction stir welding (FSW) of titanium alloys is complicated by rapid tool wear under high loads and temperatures achieved in the process. This paper studies the durability of a tool made of ZhS6U Ni-based superalloy used for friction stir processing of commercially pure titanium and the effect of the tool wear on the weld quality. The total length of the titanium weld formed by the tool without failure comprised 2755 mm. The highest wear of the tool is observed at the base of the pin, which brings about the formation of macrodefects in the processed material. The tool overheating causes an increase in the dendrite element size of ZhS6U alloy. The transfer layer contains chemical elements of this alloy, indicating that the tool wear occurs by diffusion and adhesion. As a result of processing, the tensile strength of commercially pure titanium increased by 25%.

**Keywords:** friction stir welding; grade 2 titanium alloy; ZhS6U Ni-based superalloy; microstructure; welding tool; tool wear

### **1. Introduction**

Titanium alloys have found wide applications in the aerospace, transport, and chemical industries, as well as in medicine due to their high specific strength, heat resistance, and corrosion resistance [1]. These alloys considered to be readily weldable [1], but the use of conventional fusion welding techniques on them may lead to some undesirable effects, e.g., the formation of a coarse cast structure, porosity, strains, and residual stresses. Friction stir welding (FSW), which emerged in the early 1990s, is a solid-state process [2] and is successfully used for joining various types of aluminum [3,4], magnesium [5], and copper alloys [6,7], as it has many advantages including low residual stresses. Therefore, FSW may be more suitable for joining titanium alloys than conventional fusion welding methods. In addition, the FSW is widely used for joining metallic structures including those made of dissimilar metals and are not weldable by means of fusion methods. The use of FSW for joining aluminum and titanium alloys may allow avoiding unreliable riveted joints and applying modern light materials. Furthermore, friction stir processing (FSP) may be used for surface hardening metallic materials as well as for producing bulk processed materials including composites.

A problem of the FSW and FSP techniques is the rapid tool wear under high heat and load conditions [8,9]. In view of this fact, welding tools for FSW of titanium alloys are usually made of refractory materials, including alloys based on tungsten [10–18], cobalt [19–22], and molybdenum [23–25], as well as polycrystalline cubic boron nitride [26,27]. As reported [27], the use of a polycrystalline cubic boron nitride tool for FSW of titanium alloys led to the presence of tool wear products in the weld, such as titanium boride, nitrogen, and oxygen. In addition, fabricating the cubic boron FSW tool is both a costly and tedious process. Commercially pure tungsten has good strength characteristics at elevated temperatures, but it has low fracture toughness and therefore is worn out rapidly in FSW of hard materials (steel, titanium alloys) [28]. The heat and wear resistance of a tungsten tool can be improved by alloying it with such a rare earth and scarce metal as rhenium, which makes the tool's cost even higher.

A very popular tool material for FSW of titanium alloys is cemented tungsten carbide. This material has high 1650 HV and it is slightly sensitive to abrupt temperature changes [28]. There is evidence, however, that an adhesive titanium carbide layer is formed on the surface of the FSW joints due to high affinity of titanium to carbon [29]. In addition, carbon monoxide (CO) can be formed during FSW due to high temperatures and high pressure [28]. Although, this remark is not relevant for FSW performed with argon shielding.

Some FSW tools are made of tungsten with the addition of lanthanum. It was found that the given tool material diffuses to the titanium alloy surface [15]. This fact negatively affects both the weld quality and the tool wear.

Interesting results were obtained for a cobalt-based alloy tool. It was shown that the tool wear during FSW is caused by repeated titanium alloy adhesion to and pulling out the fragments from the tool [29]. However, the durability of cobalt tools was not studied in detail. Other rarer materials were also used for FSW tools, but they have not become widespread, e.g., a molybdenum-based alloy tool [23–25]. Unfortunately, the cited authors do not describe the tool wear process and provide no data on how the use of this type of tool affects the weld quality. Another example is a zirconium diboride tool with silicon carbide additives [30], the use of which proved to be undesirable due to rapid adhesive failure. Despite the large number of studies, the available materials for the manufacture of tools are either not durable enough, do not ensure reliable weld quality, or are technologically unfeasible due to complicated processing and high cost.

The use of a polycrystalline nickel-based superalloy as a tool material for FSW of steel [31] allowed yielding positive results with respect to the tool wear despite the presence of defects in the welded joint. An iridium-containing Ni-based superalloy used for stainless steel welding showed slight wear and deformation over the welding the 3 m length seam [32]. In practice, this superalloy has been successfully applied since the 1960s for the manufacture of turbine blades operating under high temperatures and mechanical loads [33]. A recent study [34] has shown the possibility of using the Ni-based superalloy as a material for making a tool for FSW on titanium alloys. However, the tool life has not yet been considered. All the above problems also apply to friction stir processing. FSP is similar to FSW in its effects and working principle, but it is slightly simpler technologically. The objective of this study is to investigate the durability of a ZhS6U Ni-based superalloy tool and the effect of its wear on a friction stir processed commercially pure titanium for the first time.

### **2. Experimental Procedure**

### *2.1. Materials and Experimental Set-Up*

The superalloy tool life was investigated in friction stir processing (FSP) of titanium alloy workpieces under constant conditions. This process is similar to friction stir welding and differs only in the absence of a joint line between two welded parts. For convenience of description, the processed material will be called a weld. In order to avoid oxidation of titanium alloy, the processing was carried out with argon shielding. Argon was fed through a special nozzle into the processing zone. An AISI 304 stainless steel plate was used as the substrate (Figure 1). The processed titanium alloy workpieces were fixed on the welding machine table using the clamps. A rotating tool was plunged into the workpiece with an axial force and then moved forward along a straight line. Friction between the tool and the workpiece resulted in heating, plasticization and stirring of the material. The FSP was carried out at constant axial load, which is a preset parameter to maintain the metal flow both at the tool

plunging stage and during processing. The tool travel speed was not a parameter to control. The FSP tool travel started immediately after the shoulder contacted the metal surface. The tool inclination angle is 1.5◦.

**Figure 1.** Scheme of friction stir processing (**a**) and tool schematic (**b**).

The superalloy tool had a 2 mm length pin and a 20 mm diameter shoulder (Figure 1b). It was made of as-cast ZhS6U Ni-based superalloy of composition presented in Table 1. This alloy is an analogue of the Mar-M247 alloy. The tool was made using the lathe turning from a cylindrical billet.

**Table 1.** Element composition of ZhS6U alloy (wt%).

The workpieces were made of 2.5 mm thick hot-rolled sheets of grade 2 titanium alloy. The initial composition of the alloy is presented in Table 2. The FSP was carried out in the direction parallel to the rolling direction. The processing was conducted using the process parameters at which the highest weld strength was achieved as reported elsewhere [34]. The axial force on the tool during both plunging (Fpl) and processing (Fpr) was maintained at 7848 N, the traverse speed (V) was 180 mm/min, and the tool rotation rate (ω) was 950 rpm. The total weld length was more than 2.5 m. Since the laboratory FSW machine was not attached with a liquid cooling system of, unlike industrial machines, the processing was carried out in successive 100 mm length passes to avoid the tool overheating. The process was stopped after each pass to cool the tool in air. The FSP tool was not cleaned in between the passes not to interfere with the conditions created in the usual FSW process.

**Table 2.** Element composition of titanium alloy (wt%).

#### *2.2. Investigative Techniques*

Metallographic examinations were performed on both the tool and the weld materials. Samples of the processed metal were cut in a plane perpendicular to the weld. The tool material structure was examined in the cross section views. The views were obtained using polishing on abrasive papers and diamond pastes. Then, the processed titanium samples were etched in a 2% hydrofluoric acid solution for two minutes to reveal the microstructure and then washed in a 40% nitric acid solution. The tool microstructure was revealed with a reagent composed of 8 g CuSO4, 40 mL HCl, and 40 mL H2O.

Qualitative as well as quantitative analysis of the weld and tool microstructure was carried out using an Altami MET-1C metallographic microscope(Altami, St Petersburg, Russia) and a Microtrac SM 3000 scanning electron microscope (Nikkiso Co. Ltd., Tokyo, Japan) attached with an energy-dispersive X-ray microanalyzer (IXRF Systems). The worn surfaces of the tool were examined using an Olympus LEXT OLS4000 confocal microscope (Olympus, Tokyo, Japan). The tool was also photographed in the intervals between the passes to document its shape change.

Quasi-static tensile tests of the weld material were conducted in a UTC 110M-100 testing machine (Testsystems, Ivanovo, Russia) at room temperature. Flat specimens for mechanical tests were cut from stirring zone with their tensile axis coincided with the seam centerline (Figure 2). Such a test scheme was used to study the processed metal strength and thus evaluate the FSP efficiency. The base metal tensile test samples were used as reference ones. The loading rate was 1 mm/s according to the standard [35]. Microhardness numbers were obtained to evaluate the local hardness in the stir zone and in the base metal. Microhardness tester Duramin5 (Struers, Ballerup, Denmark) at 100 g load and dwell time 10 s was used.

**Figure 2.** Scheme of a tensile test specimen cut from the weld zone.

### **3. Results and Discussion**

### *3.1. Tool Wear*

The general views of the tool at different stages of the experiment are shown in optical images in Figure 3. One can see, that the tool worn surfaces are covered with an adherent layer of oxidized titanium. The edges of the pin and the shoulder are the most severely worn places during the processing. After processing the 1105 mm length, the pin shape is still conical and the depression is not observed (Figure 3b). Further processing results in gradual changing the pin shape from conical to nearly cylindrical one. A circular shape depression is gradually formed on the shoulder around the base of the pin in the course of processing.

**Figure 3.** General view of the tool before friction stir processing (FSP) (**a**), after traversing 1105 mm (**b**), 2335 mm (**c**), and 2755 mm (**d**).

Figure 4 shows typical 3D images of the tool worn surfaces before processing and after traversing 2755 mm. The surface morphology changed significantly. The adherent layers are represented by ridges generally oriented along the direction of rotation, with sparse pits formed due to adhesion and detachment of the transfer layer during processing. Despite the fact that the maximum wear of the shoulder is observed closer to the pin, there are no regular morphological features other than the fact that the shoulder roughness is higher in this area. In particular, the roughness of the pin tip surface after FSP was *Ra* = 15.6 μm, the average shoulder roughness was *Ra* = 13.6 μm, and the shoulder roughness around the pin was *Ra* = 26.2 μm (the pin and shoulder roughness before processing was *Ra* = 1.55 and *Ra* = 1.15 μm, respectively).

Figure 5 shows an SEM backscattered-electron (BSE) image of the tool worn surface after traversing 2755 mm and energy-dispersive X-ray spectroscopy (EDX) maps of Ti and W. Lighter band areas in the SEM BSE image correspond to the metal enriched with the tool material elements. It is suggested that the transfer layer was oxidized, embrittled, and detached from the tool surface when retracted from the titanium and no more protected by the argon. Therefore, this layer's fragments may be easily detached during the next FSP pass. Tungsten is also observed in deposited titanium layers, indicating the diffusion of the tool elements into the transfer layer.

The axial cross-sectional view of the tool after traversing 2755 mm is shown in Figure 6 with superimposed drawing of the initial tool geometry. The highest wear is observed in the shoulder region at the base of the pin. The initial profile height loss in this zone is as high as 1 mm. The shoulder region closer to the edge is less worn so that its height loss was 0.25 mm. The pin height remained unchanged. In our opinion, such a non-uniform wear type may be of consequences as follows. Firstly, the reduced diameter at the base of the pin can lead to insufficient heat production because of lower friction velocity and hence to poor stirring. Secondly, the pin height does not change with the shoulder wear and

therefore the tool may plunge deeper into the material. As a result, the pin can come into contact with the substrate and stick to it, or the substrate material can be introduced into the stirring zone.

**Figure 4.** Typical 3D images of the surface morphology of the tool shoulder (**a**) and pin (**c**) before processing and the shoulder (**b**) and pin (**d**) after traversing 2755 mm.

**Figure 5.** SEM BSE image of the tool friction surface after traversing 2755 mm and EDX maps of Ti and W.

**Figure 6.** Axial cross-sectional view of the tool after traversing 2755 mm with superimposed drawing of the initial tool geometry.

Costa et al. [31], despite a different pin geometry, observed a similar wear pattern of a tool from a similar Mar-M247 alloy. The authors explained the wear behavior by material overheating that causes recrystallization of the alloy, carbide formation, and oxygen diffusion, which led to cracking of the surface layers and adhesive wear. The tool damage was avoided, as reported, using the argon shielding. However, there is no explanation why the maximum wear occurs closer to the tool axis. Formation of the annular depression around the pin was explained by Liu et al. [36] using the contact melting theory when they observed a similar concavity formed by friction in the center of a coating rod after deposition. Despite the fact that the friction velocity decreases closer to the center, heat transfer is impeded in this region, as the thermal energy mainly flows from the outer radius inwards and outwards. Therefore, the rod center becomes the point where the energy is concentrated and a quasi-liquid layer is formed. That is why the wear in the center is maximum. Such an explanation seems to hold good for the case of tool wear in FSP. However, as shown [31], the formation of a depression was avoided simply by using argon shielding, which could not greatly affect the heat transfer conditions. Nevertheless, contact melting might explain the fact that the surface layer material was recrystallized.

Generally, the heat-resistant superalloy tool wear is very specific in the case of welding on titanium. For instance, in FSW with steel tool on aluminum alloys the most severe wear is experienced by the pin's end. It was shown [37] even that the pin end experienced plastic deformation. The similar wear pattern was observed on the H13 tool after welding on CuCrZr alloy [38]. FSW on Ti-6Al-4V carried out using similar configuration tools made of other materials also demonstrated different type of wear [16]. In particular, a W-La tool experienced severe plastic deformation while those made of WC-Co demonstrated brittle fracture.

An indication of the adhesive wear mechanism can be the diffusion of the tool elements into the transfer layer that remained on the friction surface after processing. Figure 7 shows an SEM image of the tool axial section and EDX maps of Ti and tool elements. The maps demonstrate that the main elements of the tool diffuse into the adherent titanium layer. The diffusion of Al is especially high. However, there is almost no titanium in the tool material. This is probably due to the fact that titanium rich layers of the tool are immediately removed during processing. Yet, the observed diffusion may be evidence of adhesive wear.

**Figure 7.** SEM image of the tool axial section and EDX maps of Ti and tool elements.

### *3.2. Structure of the Tool*

According to metallographic images of the tool and the base ZhS6U alloy (Figure 8), the structure at this scale is represented by arrays of γ-phase dendrites. The high-contrast bright objects at the array boundaries are carbides characteristic of the alloy. The shape of the arrays generally tends to be equiaxed after FSP both in the base material and in the tool. As a result of the thermal effect, the structural elements grew considerably. The size of dendritic arrays, the primary and secondary dendrite arm spacings increased by a factor of 3. The measurement results are given in Table 3. It is characteristic that the structure changed uniformly throughout the tool. The array sizes and the dendrite arm spacings did not change as well as the amount of carbides did not increase in the subsurface region. The thickness of the adherent titanium layer on the tool is in the range of 30–80 μm, which is consistent with the surface morphology measurements. The change in contrast near the worn surface is caused by over-etching due to surface tension at the edge of the specimen when it was removed from the reagent.

**Figure 8.** Metallographic images of the tool axial section after traversing 2755 mm (**a**,**c**) and the base ZhS6U alloy (**b**).

Figure 9 shows SEM BSE images of the base ZhS6U alloy structure and the tool structure. The images highlight the dendritic array boundaries decorated with eutectics and carbides of heavy metals, which mainly include Ti, W, Mo, and Nb [39]. In addition, irregular carbides and eutectics are often located in the interdendritic spaces [40]. At a finer level, the alloy structure is represented by γ'-phase cubes [41]. The measurement results for these structural elements show that the carbide volume fraction decreased slightly after FSP (Table 3), while the average carbide size increased by a factor of more than 1.5. The carbide size measurements have low accuracy because of the large size scatter in the range of 0.7–5 μm. The size change of carbides at their almost constant volume fraction is related to the coagulation of particles that occurs as dendrites grow during recrystallization. The size of γ' cubes remained unchanged as a result of the thermal effect. The material below the tool friction surface as well as below the adherent titanium layer also revealed no specific feature. Figure 9b was obtained in the same region as the optical image of Figure 8c.

**Figure 9.** Cross-sectional SEM images of the base ZhS6U alloy (**a**,**c**) and the tool material after traversing 2755 mm (**b**,**d**). Figure 9b corresponds to Figure 8c.

### *3.3. Quality of Friction Stir Processing*

The processing quality here refers to the structure of the weld and its mechanical properties. Figure 10 shows the cross-sectional metallographic images of FSPed material in different regions: at distances 5, 560, and 2755 mm from the starting point of the weld. According to metallographic analysis, the general view of the weld and the characteristic view of structural zones in all processing regions are the same for all samples. The weld structure exhibits two typical zones: stir zone (SZ) consisting of small recrystallized grains of titanium alloy, and thermomechanically affected zone (TMAZ). The rest is the base material. The heat affected zone, which is usually located between the base metal and TMAZ, is often absent or not detected in titanium alloys. The TMAZ in FSPed titanium alloy is narrow as compared to that of obtained on materials, with the average thickness of about 0.25 mm. The geometry of the structural zones is generally symmetrical about the weld axis. Their shapes in different processing regions differ slightly. The SZ area gradually decreases with the tool wear: from 11.11 mm2 in the first sample to 8.03 mm2 in the last one. This occurs because the pin diameter is reduced as the tool is worn.