# Heterogeneous nucleation and orientation relationships of icosahedral phase with TiB2 inoculants

Blaž Leskovar, Sašo Šturm, Kemal Delijić, Julian Ledieu, Matejka Podlogar, Sandra Drev, Iztok Naglič, Boštjan Markoli

► To cite this version:

Blaž Leskovar, Sašo Šturm, Kemal Delijić, Julian Ledieu, Matejka Podlogar, et al.. Heterogeneous nucleation and orientation relationships of icosahedral phase with TiB2 inoculants. Journal of Alloys and Compounds, In press, 968, pp.172195. 10.1016/j.jallcom.2023.172195. hal-04243256

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Submitted on 16 Oct 2023

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Research article

# Heterogeneous nucleation and orientation relationships of icosahedral phase with TiB₂ inoculants

<figure><img src="image_2.png" /></figure>

Blaž Leskovar<sup>a, *</sup>, Sašo Šturm<sup>b, f, g</sup>, Kemal Delijić<sup>c</sup>, Julian Ledieu<sup>d</sup>, Matejka Podlogar<sup>b</sup>, Sandra Drev<sup>e</sup>, Iztok Naglič<sup>a</sup>, Boštjan Markoli<sup>a</sup>

## ARTICLE INFO

### Keywords:

Aluminium alloys

Casting

Heterogeneous nucleation

Orientation relationship

Icosahedral quasicrystals

## ABSTRACT

This study confirms that crystalline TiB₂ particles can act as a potent site for the heterogeneous nucleation of the metastable icosahedral quasicrystalline phase (IQC) formed in cast Al-Mn-Si-Cu-Mg alloys. The results show that a primary metastable IQC phase nucleates on facets of TiB₂ particles. Electron backscatter diffraction patterns and selected-area electron diffraction patterns confirmed that at least five orientation relationships (ORs) exist between the stable TiB₂ and the metastable IQC phase. The IQC adopts different ORs with the crystalline TiB₂ particles to form low energy interfaces by matching close-packed planes at the interfaces. In the IQC phase, the close-packed planes are the 5-fold and 2-fold. Microstructural analyses show that inoculation has a noticeable influence on the size and distribution of the primary IQC phase, which in turn has a noticeable influence on the compressive properties of the investigated alloys. Due to their exceptional mechanical properties, these alloys could become a prime candidate for lightweight applications, especially in the automotive and aerospace industries.

## 1. Introduction

Icosahedral quasicrystals (IQCs), first discovered in the alloy system Al-Mn, were thermodynamically metastable and formed only when the melt was rapidly cooled (10⁶ K/s). When the alloy was annealed, decomposition occurred in the thermodynamically more stable crystal-line phase known as the approximant [1,2]. Soon after the discovery of the metastable IQC phase, stable IQCs were also synthesized in various alloys at very specific compositions by slow cooling of the alloy melt [3, 4]. Continuous progress in materials science offers great potential also for innovation of alloys reinforced with IQC phases, especially for lightweight applications in the automotive and aerospace industries. From the field of microstructure engineering by solidification processing, composite-like materials containing IQC are the subject of various application experiments in terms of improved mechanical performance, whether by blending with polymers or light metals or by in-situ

reinforcement produced by nanoprecipitation of particles [5,6]. A large number of Al-based alloys reinforced with metastable IQCs are known. Such alloys, which can be produced by rapid solidification techniques (10³–10⁶ K/s), exhibit excellent thermal resistance, high strength at room temperature, high strength and hardness at high temperatures, excellent ductility and high specific strength [7–13]. In some aspects, they are even better than commercial high-strength Al-alloys, which are used for a variety of applications in the marine, aerospace and automotive industries [5]. The main drawback in the production of composite-like materials containing metastable IQC phases is that the volume fraction, distribution, and morphology of the IQC phase are highly dependent on the composition and process parameters [14–16]. To obtain a fine, uniform, equiaxed structure of the primary IQC phase, precise control of nucleation and growth is of great importance. Inoculation of melts to suppress columnar solidification and to obtain fine equiaxed grain structures in the cast solid is now a common industrial

<footnote><sup>a</sup> University of Ljubljana, Faculty of Natural Sciences and Engineering, Department of Materials and Metallurgy, Aškerčeva cesta 12, Ljubljana 1000, Slovenia</footnote>

<footnote><sup>b</sup> Jožef Stefan Institute, Department for Nanostructured Materials, Jamova cesta 39, Ljubljana 1000, Slovenia</footnote>

<footnote><sup>c</sup> University of Montenegro, Faculty of Metallurgy and Technology, Cetinjski put b.b., Podgorica 81000, Montenegro</footnote>

<footnote><sup>d</sup> Université de Lorraine, CNRS, Institut Jean Lamour, Nancy F-54000, France</footnote>

<footnote><sup>e</sup> Jožef Stefan Institute, Center for Electron Microscopy and Microanalysis, Jamova cesta 39, Ljubljana 1000, Slovenia</footnote>

<footnote><sup>f</sup> International Postgraduate School Jožef Stefan, Jamova cesta 39, Ljubljana 1000, Slovenia</footnote>

<footnote><sup>g</sup> University of Ljubljana, Faculty of Natural Sciences and Engineering, Department of Geology, Aškerčeva cesta 12, Ljubljana 1000, Slovenia</footnote>

<footnote>* Corresponding author.
E-mail address: blaz.leskovar@ntf.uni-lj.si (B. Leskovar).</footnote>
practice, but it has never been studied for the primary metastable IQCs phase. During inoculation, a master alloy is added to the melt to increase its solute content and to introduce stable particles that can serve as substrate for grain growth. It is known that for any phase to act as an effective substrate there must be a good lattice match with the phase that is being nucleated. The orientation relationship between the phases occurs in order to minimize the free energy of the interface. Low energy interfaces usually form between the close-packed planes of the substrate and the growing phase [17–19].

A fundamental experiment based on the lattice and structural similarity between the metastable IQC-AlMnSi and the stable IQC-AlCuFe quasicrystals was performed in a previous study by Leskovar et al. [20]. The results show that a continuous thin layer of the metastable IQC-AlMnSi phase was formed on the stable IQC-AlCuFe particle during the solidification of the alloy Al-Mn-Si. This clearly indicates that particles of stable IQC-AlCuFe can act as an efficient substrate surface on which the metastable IQC-AlMnSi can grow. The main shortcoming of the stable IQC-AlCuFe phase, which was used as a substrate for the nucleation of the primary metastable IQC-AlMnSi phase, is that it decomposes rapidly by being held in the Al-based melt [20].

More sustainable inoculants for nucleation of the primary metastable IQC phase can thus present more stable crystalline phases in the Al-based melt. It is reported in the literature that a specific orientation relationship (OR) between quasicrystalline and closely related crystalline structures (approximants) was observed shortly after the discovery of quasicrystals [21–23]. Review articles present the results of the growth of atomic overlayers of various elements on different quasicrystalline substrates, with the main goal of obtaining pseudomorphic quasicrystalline films of single elements [24]. Only a limited number of elements have yielded films with quasiperiodic structure, while most elements develop periodic films that exhibit an OR with the substrate. Other interesting phenomena associated with growth have also been observed, including intermixing, surface alloying, twinning of nanocrystals, and heterogeneous nucleation at quasilattice sites [24–27]. The orientation relations of quasicrystal-crystal interfaces in bulk materials were presented by Singh and Tsai [28]. They found that the crystals form various orientational relationships with the quasicrystals in order to form low-energy interfaces by matching close packed planes across the interfaces.

Fleury et al., [29], were among the first to report the heterogeneous nucleation of the metastable IQC phase on the crystal phase. Microstructural and thermal analyses indicated that the IQC phase nucleated heterogeneously from the cubic AlMnBe₄ phase by a peritectic reaction, but with a slight misorientation [29]. The AlMnBe₄ phase was not added to the melt as substrate but was already present in the melt, due to its high melting point (900 °C), as determined by thermal analyses.

These results clearly confirm that the crystalline phases can serve as a substrate for the heterogeneous nucleation of the primary metastable IQC phase in the Al-Mn-based alloys. A brief search for the possible crystalline inoculants, based on the lattice spacing of the close-packed planes and the stability in the Al-based melt, revealed that the TiB₂ particles which are commonly used for α₉₁ nucleation in commercial Alloys can also be used as substrate for the nucleation of the primary metastable IQC phase. The difference between the close-packed planes (5-fold and 2-fold) of primary IQC [30] and TiB₂ (a=0.3029 nm and c=0.3220 nm) [31,32] is in the range of 0.5–6 %.

Therefore, the aim of this study was to elucidate and evaluate the role of the inoculation of cast Al-Mn-Si-Cu-Mg alloy with TiB₂ particles on the formation of metastable IQCs and its impact on the mechanical properties of cast alloy.

## 2. Materials and methods

To evaluate the effect of inoculation, two Al-Mn-Si-Cu-Mg (BIO1) alloys were prepared in a chamber furnace by melting pure aluminium (99.8 wt%), copper (99.9 wt%), manganese (99.9 wt%), magnesium

(99.9 wt%), and silicon (99.99 wt%) at 880 °C. One alloy was cast with and the other without the addition of AlTi3B1 master alloy (inoculants – TiB₂ particles). The BIO1 alloy was selected due to its ability to form metastable IQCs throughout the casting volume at cooling rates of about 600 K/s, as previously reported [15].

The commercial master alloy AlTi3B1 was added to the melt five minutes before casting into the copper mould. The amount of AlTi3B1 alloy was calculated according to the total amount of titanium in the alloy, which was 0.04 wt% Ti.

In the AlTi3B1 master alloy, most of the titanium (≈ 74 %) is combined with boron to form a TiB₂ phase. The remainder of the titanium is bound to the Al₃Ti phase, and a small portion of the titanium is soluble in α₉₁, as is also known from the literature [17–19] and the Al-Ti-B phase diagram [33,34].

The chemical composition of the synthesized alloys (BIO1) was determined by the EDS, Table 1. Titanium and boron are not included in Table 1 because their corresponding experimental uncertainties are much higher than the measured content.

The secondary electron image (SEI) of the etched AlTi3B1 master alloy showed that the TiB₂ phase has a faceted hexagonal platelet morphology with {0001} planes as basal facets and {0110} and {0111} planes as prism facets, Fig. 1. This morphology is similar to the theoretical growth morphology calculated from the binding energies of the crystal facets [31,32]. The average diameter of TiB₂ particles is estimated to be around 1 µm.

The casting mould was made from a pure copper block with dimensions 100 × 100 × 120 mm and a cylindrical casting cavity with a diameter of 5.5 mm, Fig. 2a. The sample for electron microscopy and microanalysis was cut 5 mm from the tip of the 5.5 mm diameter round casting. Two samples for compression testing, 13 mm in length, were cut later, as shown in Fig. 2b. The sample for microstructure analysis was prepared by grinding and final polishing with colloidal silica of grain size 0.05 µm. The homogeneity and size of the primary IQC phase were checked in the area close to the contact surface and in the central area of the investigated samples, Fig. 2c.

Microstructure was characterised using a light optical microscope (LOM), ZEISS Axio Imager A1m, equipped with an AxioCam ICc 3 digital camera (3.3 million pixels) and AxioVision software. The primary IQC phase area fraction was determined by software processing of LOM images acquired over the entire surface of the sample at a magnification of 100 ×. To cover the sample cross-section, 24 microstructures were analysed. In order to obtain more reliable statistical certainty, cross sections of six alloys (three BIO1 and three BIO1-TiB2) were analysed at a distance of 5 mm from the tip of the casting (Fig. 2). Since the area fraction of symmetrical phases is equal to the volume fraction, we will refer further to the volume fraction [35]. A JEOL JSM-7600 F field emission scanning electron microscope was used for microstructural characterization. Elemental and micro-crystallographic analyses were performed using the Oxford Instruments INCA Microanalysis Suite, by energy dispersive X-ray spectroscopy (EDS) with an X-Max 20 SDD detector and by electron backscatter diffraction (EBSD) with a Nordlys detector and CHANNEL5 software. The crystallinity and orientation of the phases were investigated using the EBSD technique at 20 kV, while the composition of the alloys and phases was determined using EDS at 20 and 5 kV. Rendering of observation of the EBSD with respect to a sample inserted into a SEM chamber is shown in Fig. 3a. The EBSD system used was checked according to the procedure of Britton et al. [36] and it was confirmed that we used the right-handed set with 70°

<table><caption>**Table 1**
Average bulk chemical composition of the prepared alloys as established by EDS in at% based on 12 measurements for each alloy.</caption><thead><tr><th>Alloy</th><th>Al</th><th>Mn</th><th>Si</th><th>Cu</th><th>Mg</th></tr></thead><tbody><tr><td>BIO1</td><td>89.8 ± 0.3</td><td>4.1 ± 0.3</td><td>1.5 ± 0.1</td><td>1.4 ± 0.2</td><td>3.2 ± 0.2</td></tr><tr><td>BIO1-TiB2</td><td>89.8 ± 0.3</td><td>4.0 ± 0.2</td><td>1.5 ± 0.2</td><td>1.5 ± 0.1</td><td>3.2 ± 0.2</td></tr></tbody></table>
<figure><img src="image_3.png" /><figcaption>**Fig. 1.** Deep-etched surface of the AlTi3B1 master alloy with the TiB₂ particles. a) TiB₂ particles showing hexagonal prism morphology. b) Growth morphology of TiB₂ particles with marked facets planes.</figcaption></figure>

<figure><img src="image_4.png" /><figcaption>**Fig. 2.** Sampling procedure. a) Copper mould with the cast alloy. b) Sectional diagram of a cast specimen with a diameter of 5.5 mm for LOM, SEM and compression tests. c) Sketch of the sample for LOM and SEM analysis with the marked sampling region and the edge and central area, based on which we inferred the efficiency of the inoculants.</figcaption></figure>

<figure><img src="image_5.png" /><figcaption>**Fig. 3.** Schematic representation of EBSD data acquisition. a) The orientation of the sample reference frame, RD-TD-ND and Xs-Ys-Zs; and the EBSD pattern reference frame [36]. b) The IQC reference frame, where the natural basis vectors of Cartesian coordinates are those along three mutually orthogonal 2-fold axes.</figcaption></figure>

sample tilt and with 10° camera tilt. The approach of Tanaka et al.[37] was used to determine the Euler angles of the IQC phase, where the natural basis vectors of the Cartesian coordinates for IQC were those

along three mutually orthogonal 2-fold axes, as shown in Fig. 3b.

A FEI HeliosNanolab 650 focused ion beam (FIB) was used to prepare the sample for transmission electron microscopy (TEM). Gallium ions
were used in the preparation process, which began with high-energy milling (30 keV) and ended with low-energy final polishing (2 keV) of the sample. A Jeol JEM-2100 transmission electron microscope equipped with a Gatan CCD camera ORIUS was used for microstructure and electron diffraction characterization. Cylindrical specimens with a diameter of 5.5 mm and a height of 13 mm were used for the compression tests (Fig. 2b). The tests were performed on the universal uniaxial device VEB Thüringen Industriewerk Betrieb des WMK Fritz Heckert "Rauenstein" - FPZ 100/1. The mechanical properties of the tested specimens (uniaxial compression) were determined based on the curve of displacement and compressive force. The maximum compressive strength (MPa) was calculated based on the local maximum of the compressive force and divided by the initial cross-section. Compressive fracture strain was calculated based on the initial and final heights of the specimen and expressed as a percentage (%). The compressive yield strength was calculated based on the compressive force at 0.2 % deformation and divided by the initial cross-section of the specimen.

## 3. Results

### 3.1. Microstructures

The LOM figures of the cast BIO1 and BIO1-TiB2 alloys in the area near the contact surface (Fig. 4a and c) and in the central area of the casting (Fig. 4b and d) are shown together in Fig. 4. A primary icosahedral phase with a dendritic morphology, often exhibiting a five-branch feature, is present throughout the cross-section of the castings. The primary IQC dendritic phases in the central region of the samples (Fig. 4b and d) are larger, with longer dendritic branches compared to those found near the contact surface of the casting (Fig. 4a and 4c) for

both investigated alloys. The primary branches in the central region of the BIO1 alloy casting (Fig. 4b) are about 50 µm long, while those in the BIO1-TiB2 alloy casting (Fig. 4d) are only about 30 µm long.

From the LOM figures, one can see that with the addition of inoculants, the IQC phase is more uniformly distributed over the sample cross-section. In the area near the contact surface, the IQC phase in the BIO1-TiB2 alloy has larger dendritic branches, while they are shorter in the centre of the cast sample in comparison to BIO1 alloy.

The area fraction of the primary IQC phase was determined from LOM images taken at 100 × magnification over the entire surface of the sample. It was assumed that the area fraction of the IQC phase is equal to the volume fraction, as explained in the Materials and Methods section. To cover the entire sample cross-section (from the area near the contact surface to the central area (Fig. 2a)), 24 images of microstructures had to be analysed. In order to obtain more reliable statistical information, whole cross-sections of three samples of alloy BIO1 and three samples of alloy BIO1-TiB2 were analysed, resulting in a total of 72 individual images of microstructure. The average volume fraction of the primary icosahedral phase is 18.1 ± 0.6% for the BIO1 alloy and 18.6 ± 0.4% for the BIO1-TiB2 alloy. Fig. 5 shows the distribution of the volume fraction of the IQC phase for each of 72 microstructures. Each volume fraction was assigned to a different volume class (eight classes altogether), spanning from below 6% all the way up to 24% and more to obtain volume fraction distribution for IQC phase over the cross-section.

In the BIO1 alloy, there is a large difference between the area near the contact surface and the central region. The near-contact area is dominated by microstructures with lower content of primary IQC phase, which can be clearly seen in Fig. 5, where an IQC volume fraction between 6% and 9% appears for the first maximum (Max 1) of the number of microstructures. There is also a higher number of microstructures

<figure><img src="image_6.png" /><figcaption>**Fig. 4.** LOM microstructures of the investigated alloys. a) Area near the contact surface and, b) central region of the BIO1 alloy in the as-cast state. c) area near the contact surface and, d) central region of the BIO1-TiB2 alloy in the as-cast state. The LOMs in the insets were made with higher magnification to show the IQC morphology in more detail.</figcaption></figure>
<figure><img src="image_7.png" /><figcaption>**Fig. 5.** The volume fraction of the primary IQC phase over three sample cross sections of the BIO1 (grey) and BIO1-TiB2 (blue) alloy, divided into eight volume classes. This results in 72 individual images of microstructures analysed for each alloy. The maxima and average volume fraction class for both alloys are marked in the image.</figcaption></figure>

with volume fraction above 21 %, indicating the second maximum (Max 2). These microstructures mainly represent the central region of the sample. The reason for the higher portion of IQC phase in the central region is due to the directional solidification and growth of the primary IQC phase from the cold Cu-mould towards the centre of the sample. It is assumed that part of the IQC phases formed in the area near the contact surface and migrated toward the centre of the sample with the solidification front.

In contrast, the amount of IQC phase in the BIO1-TiB2 alloy is more uniformly distributed (Gaussian) across the cross section. The majority (Gaussian maximum) of the microstructures belong to the class of average content of primary IQC phase (18.6 ± 0.4 %) in the range of 18–21 %, which we believe is related to the increased number of nucleation sites due to TiB₂ addition.

Analysis of the microstructure by electron microscope showed that TiB₂ particles (lighter white colour) were located in the centre of the primary IQC phase (see Fig. 6). These particles clearly indicate that the crystalline phase was already present in the melt before the primary IQC phase formed. The formation of IQC around the crystalline phase could be explained by two interrelated aspects. First, the crystal phase served

as a heterogeneous nucleation site for the growth of the primary phase and second, there must have been a structural similarity in defined crystallographic planes that have the lowest interfacial free energy in order to allow growth with respect to the specific crystal/quasicrystal direction. The region with TiB₂ particles in the middle of the primary IQC phase was extracted from the sample using FIB and prepared for TEM study.

## 3.2. Orientation relationship

The orientation relationships (ORs) between the hexagonal (h) TiB₂ particles as a nucleation site and the primary IQC phase were investigated using the EBSD and TEM diffraction methods. Based on the EBSD images, five different ORs were identified and designated as (hORs). For a better understanding of the hORs, the index (100000) stands for the 5-fold symmetry plane and (010100) for the 2-fold symmetry plane of the IQC phase.

The SEI images (insets in Fig. 7a–j) show that TiB₂ particles are located within the central part of the IQC phase. The orientation of the TiB₂ particles with respect to the sample surface is represented by Euler angles calculated using CHANNEL5 software, while the orientation of the IQC phase was calculated manually using the approach of Tanaka et al. [37].

The first orientation relationship (hOR1) is shown in Fig. 7a and b. The diffraction patterns of TiB₂ and the IQC phase show that the planes (0001) in TiB₂ and (100000) in the IQC phase are parallel. Moreover, one can see that the planes (1010) in TiB₂ and (001100) in IQC are also parallel. The orientation of the TiB₂ particle and the IQC phase with respect to the sample surface represented by the Euler angles are shown in Fig. 7a and b.

The second orientation relationship (hOR2) found between the phases is presented in Fig. 7c and d. The diffraction patterns of TiB₂ and the IQC phase show that the directions [0001] for TiB₂ and the [100000] for IQC phase are parallel. In addition, it is also apparent that the planes (1210) in TiB₂ and (010100) in IQC are also parallel, thus directions [1210] for TiB₂ and [010100] for IQC must also be parallel. The orientation of the TiB₂ particle and the IQC phase with respect to the sample surface represented by the Euler angles are shown in Fig. 7c and d.

The third orientation relationship (hOR3) found between phases is presented in Fig. 7e and f. The diffraction patterns of TiB₂ and the IQC

<figure><img src="image_8.png" /><figcaption>**Fig. 6.** An SEM image of the primary IQC phases with two TiB₂ particles in the middle; the inset shows a FIB lamella extracted from the centre of IQC phase including TiB₂ particle. The circles show the area of interest.</figcaption></figure>
EBSD diffraction patterns

<figure><img src="image_9.png" /></figure>

Euler angles

<figure><img src="image_10.png" /></figure>

**Fig. 7.** EBSD diffraction patterns and corresponding 3D orientation with Euler angles of TiB₂ phase (Figures (a), (c), (e), (g), and (i)) and IQC phase (Figures (b), (d), (f), (h), and (j)). In the inset of the EBSD diffraction patterns, there is an SEI image of the TiB₂ particle in the central part of the IQC phase with a labelled location (white cross) of the EBSD analyses. Parallel planes and directions between phases are drawn. The Euler angles of the IQC phase were calculated according to the procedure of Tanaka et al. [37]. The orientation relationship between TiB₂ and IQC is represented as follows: hOR1 in Figures (a) and (b), hOR2 in Figures (c) and (d), hOR3 in Figures (e) and (f), hOR4 in Figures (g) and (h), and hOR5 in Figures (i) and (j). In the EBSD diffraction patterns, the characteristic 5-fold (white five-pointed star), 3-fold (white triangle), and 2-fold (white ellipse) axes are marked.

phase show that the directions [11̄20] for TiB₂ and [000100] for the IQC phase are parallel. It is also seen that the planes (11̄00) in TiB₂ and (011000) in IQC are parallel, thus directions [11̄00] for TiB₂ and [011000] for IQC must also be parallel. The orientation of the TiB₂ particle and the IQC phase with respect to the sample surface represented by the Euler angles are shown in Fig. 7e and f.

The fourth orientation relationship (hOR4) is shown in Fig. 7g and h. The diffraction patterns of TiB₂ and the IQC phase show that the planes (11̄20) in TiB₂ and (01̄0010) in the IQC phase, along with the planes (11̄22) in TiB₂ and (01̄0100) in the IQC phase are parallel. Moreover, one can see that the directions [11̄00] in TiB₂ and [100000] in IQC phase are also parallel. The orientation of the TiB₂ particle and the IQC phase with respect to the sample surface represented by the Euler angles are shown in Fig. 7g and h.

The fifth orientation relationship (hOR5) found between the phases is presented in Fig. 7i and j. The diffraction patterns of TiB₂ and the IQC phase show that the planes (11̄00) in TiB₂ and the (001001) in IQC phase

and planes (0001) in TiB₂ and (000110) in IQC phase are parallel. Based on the obtained OR, we conclude that directions [11̄00] for the TiB₂ and [001001] for the IQC phase; and directions [0001] for the TiB₂ and [000110] for the IQC phase must also be parallel. From Fig. 7i and j, one can see that directions [11̄20] for the TiB₂ and [110000] for the IQC phase are also parallel. The orientation of the TiB₂ particle and the IQC phase with respect to the sample surface represented by the Euler angles are shown in Fig. 7i and j.

The stereographic projection of the hOR1 is shown in Fig. 8a and b. From the orientation relation obtained, one can see that the [0001] and [0001] directions for TiB₂ and 5-fold symmetry axis for IQC phase and that the direction [1010] for TiB₂ and 2-fold symmetry axis for IQC phase are parallel.

The hOR2 is shown by stereographic projection of the phases in Fig. 8c and d. From the obtained orientation relation, one can see that the [0001] and [0001] directions for TiB₂ and 5-fold symmetry axis for IQC phase and that the direction [12̄10] for TiB₂ and 2-fold symmetry
<figure><img src="image_11.png" /><figcaption>**Fig. 8.** Stereographic projections show the orientation relation between TiB₂ and IQC phases with marked matching directions for the TiB₂ phase (Figures (a), (c), (e), (g) and (i)), and the IQC phase (Figures (b), (d), (f), (h) and (j)). The circled directions in both stereographic projections are parallel (or with angle mismatches less than 1°), whereas dotted circles show directions with angle mismatches from 1° to 5°. The orientation relationship between TiB₂ and IQC is represented as: hOR1 in Figures (a) and (b), hOR2 in Figures (c) and (d), hOR3 in Figures (e) and (f), hOR4 in Figures (g) and (h), and hOR5 in Figures (i) and (j).</figcaption></figure>

axis for IQC phase are parallel. Moreover, the direction [1011] for TiB₂ and 2-fold symmetry axis for IQC phase have almost epitaxial correspondence. The angle between the [0001] and [1011] directions in TiB₂ is 58.34°, while the angle between the direction of 5-fold and 2-fold symmetry axis (over 3-fold symmetry axis) in IQC is 58.29°. Thus, the

difference is only 0.05°.

The hOR3 is shown by stereographic projection in Fig. 8e and f. From the OR obtained, one can see that the directions [11̄20] for TiB₂ and 5-fold symmetry axis for IQC phase and the directions [1̄100] for TiB₂ and 2-fold symmetry axis for IQC phase are parallel. A small mismatch of 4.3° between the [1̄101] and [1̄1001] directions for TiB₂ and neighbouring 2-fold symmetry axes in the IQC phase can also be seen.

The hOR4 is shown by stereographic projection of the phases in Fig. 8g and h. From the OR obtained, one can see that the directions [11̄20] for TiB₂ and 2-fold symmetry axis for IQC phase and that the [1̄100] and [1̄100] directions for TiB₂ and 5-fold symmetry axes for IQC phase are parallel. However, if we use the [11̄20] direction as the zone direction for the stereographic representation of the OR and the [1̄100] direction for TiB₂ as the epitaxial matching with the 5-fold direction for the IQC phase, we obtain a mismatch of 0.5° between the directions [11̄22] for TiB₂ and the 2-fold symmetry axis for the IQC phase and the directions [11̄22] for TiB₂ and the 2-fold symmetry axis for the IQC phase.

The hOR5 is shown by stereographic projection of the phases in Fig. 8i and j. From the OR obtained, one can see that in addition to the epitaxial matching between the TiB₂ and IQC phase directions explained by the EBSD diffraction patterns, there is also a small mismatch between other main directions in the phases. A mismatch of 1.7° between the directions [011̄0] for TiB₂ and the 5-fold symmetry axis for the IQC phase and [101̄0] for TiB₂ and the 5-fold symmetry axis for the IQC phase, calculated from the stereographic zone directions of the phases ([1100] for TiB₂ and 2-fold axis for IQC), can also be seen. Here, an almost epitaxial matching between the directions [11̄01] in TiB₂ and the 5-fold symmetry axis in IQC phase as well as [1̄101] for TiB₂ and the 5-fold symmetry axis for IQC phase is seen, with mismatch of 0.06°.

The interface region between the TiB₂ particle and the IQC phase was further investigated using TEM to confirm the OR obtained by the EBSD method (Fig. 9). Fig. 9a shows the bright-field TEM image of the IQC phase with TiB₂ particle in the centre and surrounding αAl-matrix. The interface between TiB₂ particle and IQC phase is free from any other phases, showing that the interface is indeed completely coherent. The IQC phase nucleated on the pure surface of the TiB₂ particle and grew into the surrounding melt. Fig. 9b shows the selected-area electron diffraction (SAED) patterns of the TiB₂ particle and the IQC phase, both acquired at the same orientation of the sample. The SAED patterns indicate that the direction of the incident beam is parallel (zone axis) to the 2-fold symmetry axis of IQC and [1̄100] of the TiB₂ phase. One can clearly see that the 5-fold planes (200000) of the IQC phase and the (0003) planes of the TiB₂ phase are also parallel. This orientation relation is consistent with the hOR1 obtained by the EBSD technique.

The interplanar spacings for the (200000) planes for IQC and (0003) planes for TiB₂ were measured using individually taken SAED patterns (i.e., slightly away from the particle). The measured interplanar spacings were 0.1086 ± 0.01 nm for the IQC phase and 0.1078 ± 0.01 nm for the TiB₂. The measured values are in good agreement with the reported (200000) interplanar spacings, which are 0.1085 nm for the metastable Al-Mn IQC phase [30] and 0.1073 nm for the crystalline TiB₂ [32]. The misfit calculated from the measured interplanar spacings for the (200000) planes for the IQC and (0003) for the TiB₂ phase is 0.737 %, which agrees well with the misfit calculated from the published interplanar spacings, i.e., 1.106 %. The results presented here support the hypothesis that the TiB₂ particles can serve as a substrate for the nucleation of metastable IQC phase and can thus be used to induce controlled formation of IQC particles in Al-based alloys. Table 2 shows epitaxial orientation matching of hexagonal TiB₂ phase with IQC phase in the orientation relationships obtained in this study.

Among the 24 analysed locations, six hOR1, six hOR2, five hOR3, two hOR4 and five hOR5 were found. The reason for the lower amount of hOR4 could be that here the growth is less energetically favoured than
<figure><img src="image_12.png" /><figcaption>**Fig. 9.** Bright-field image and Selected-area electron diffraction (SAED) pattern of TEM sample. a) IQC phase with TiB₂ particle in the centre, and surrounding α-Al-matrix. b) SAED pattern of the interface between TiB₂ and IQC phase. The white circle in Fig. 9a denotes the location where the SAED pattern was taken, while the white circles in Fig. 9b denote peaks representing TiB₂ planes. The matching planes between TiB₂ and IQC are marked with white arrows. This orientation relationship corresponds to hOR1, which was determined using the EBSD technique in this work.</figcaption></figure>

<table><caption>**Table 2**
Epitaxial orientation matching between hexagonal TiB₂ and IQC phases.</caption><thead><tr><th>hOR</th><th>Icosahedral planes</th><th>Corresponding TiB<sub>2</sub> planes</th><th>Plane Lattice Mismatch</th></tr></thead><tbody><tr><td rowspan="2">hOR1</td><td>Indices / Symmetry {200000} / 5-fold</td><td>Spacing (Å) 1.085</td><td>Indices {0003} 1.073</td><td>(%) 1.106</td></tr><tr><td>{001001} / 2-fold</td><td>Spacing (Å) 1.275</td><td>{2020} 1.312</td><td>(%) -2.902</td></tr><tr><td rowspan="2">hOR2</td><td>{200000} / 5-fold</td><td>Spacing (Å) 1.085</td><td>Indices {0003} 1.073</td><td>(%) 1.106</td></tr><tr><td>{330011} / 2-fold</td><td>Spacing (Å) 1.496</td><td>{1210} 1.515</td><td>(%) -1.270</td></tr><tr><td rowspan="2">hOR3</td><td>{622222} / 5-fold</td><td>Spacing (Å) 1.428</td><td>{1120} 1.515</td><td>(%) -6.092</td></tr><tr><td>{101000} / 2-fold</td><td>Spacing (Å) 1.275</td><td>{2200} 1.312</td><td>(%) -2.902</td></tr><tr><td rowspan="2">hOR4</td><td>{330011} / 2-fold</td><td>Spacing (Å) 1.496</td><td>{1120} 1.515</td><td>(%) -1.270</td></tr><tr><td>{311111} / 5-fold</td><td>Spacing (Å) 2.856</td><td>{1100} 2.623</td><td>(%) 8.158</td></tr><tr><td rowspan="3">hOR5</td><td>{330011} / 2-fold</td><td>Spacing (Å) 1.496</td><td>{1120} 1.515</td><td>(%) -1.270</td></tr><tr><td>{101000} / 2-fold</td><td>Spacing (Å) 1.275</td><td>{2200} 1.312</td><td>(%) -2.902</td></tr><tr><td>{220000} / 2-fold</td><td>Spacing (Å) 1.037</td><td>{0003} 1.073</td><td>(%) -3.472</td></tr></tbody></table>

in the others. It is clear from Table 2 that hOR4 has the highest mismatch for the second plane.

### 3.3. Compressive properties

As pointed out in the introduction, the controlled formation, distribution, and size of IQC phase could improve mechanical properties of Al-Mn-based alloys and make them a preferred choice for structural applications. The compression tests, based on five cylindrical specimens of each alloy, show that the compressive fracture strain increased significantly with the addition of TiB₂ particles, while the compressive yield strength and compressive strength did not change much and remained within the statistical error range (Graph 1). It is also evident that the size of the IQC phase affects the properties. Shorter dendritic branches of the primary IQC phase (Fig. 4) have a significant effect on reducing the notch effect and thus on the higher ductility of these alloys.

## 4. Discussion

Previous studies have found that an orientation relationship exists between hexagonal phases structurally related to quasicrystals (approximants) and the icosahedral phase in which an icosahedral 2-fold axis occurs along the hexagonal axis in a symmetrical fashion, with another two 2-fold axes along 2-fold axes of the hexagonal phase [21–23].

We have also confirmed that TiB₂ shows this OR with the IQC phase,

<figure><img src="image_13.png" /><figcaption>**Graph 1.** Compressive yield strength, compressive fracture strain and compressive strength for BIO1 and BIO1-TiB₂ alloys.</figcaption></figure>

but since the TiB₂ has a simple structure not related to the icosahedral phase it shows other ORs too.

A comprehensive comparison of all ORs found in our study (Fig. 10a) and in previous studies (Fig. 10b) is presented in Fig. 10.

In hOR1, one 5-fold plane exists parallel to the hexagonal basal plane. The set of 2-fold planes match the set of prismatic planes {2020}. The hOR2 is a variant of hOR1, related by a rotation of 90° around the coincident axis. In this, a set of 2-fold planes match {1210} planes. In hOR3 and hOR4 {τ10}, zone axis occurs along the hexagonal axis. In hOR3, one of the 2-fold planes matches a prismatic plane, and one 5-fold plane matches a {1120} plane, whereas in hOR4 one of the 5-fold planes matches a prismatic plane, and one 2-fold plane matches a {1120} plane. The hOR3 is a variant of hOR4, related by a rotation of 90° around the coincident axis. The hOR5 is similar to the OR observed between the IQC and the quasicrystal-related (approximants) hexagonal phases. Where the icosahedral 2-fold axis occurs along the hexagonal axis, a set of prismatic planes match icosahedral 2-fold planes and a set of {1120} planes match other icosahedral 2-fold planes.

Similar findings were presented by Sing and Tsai [28] where interfaces of icosahedral phase particles in a simple hexagonal magnesium matrix were characterized. Magnesium is hexagonal, with a nearly ideal c/a axial ratio. They found five ORs of IQC phase in a magnesium matrix. Of these, MgOR1 is similar to the OR observed between the IQC phase and the quasicrystal-related (approximants) hexagonal phases and is the same as hOR5 in this study. MgOR2 is a variant of MgOR1, related by a rotation of 90° around the coincident axis. In this, two 5-fold planes match {1120} planes, while one 2-fold plane matches a set of prismatic planes. This OR was not found in our study, but most probably exists. MgOR3 is equivalent to hOR1 in this study where one 5-fold plane exists parallel to the hexagonal basal plane. One set of 2-fold planes match a set of prismatic planes. In MgOR4 a {τ10}, zone axis occurs along the hexagonal axis. In this OR, one of the 5-fold planes matches a prismatic
<figure><img src="image_14.png" /><figcaption>**Fig. 10.** The five ORs of IQC and TiB₂ obtained in this study and compared to other hexagonal ORs. a) five ORs of IQC phase in this study, b) modified ORs from Singh and Tsai [28]. The hexagons represent the orientation of the TiB₂ and magnesium lattice along its hexagonal axis and the pentagonal dodecahedron shows the corresponding orientation of the IQC phase.</figcaption></figure>

plane, and one 2-fold plane matches a {1120} plane. This OR is equal to hOR4 in this study. In MgOR5, an icosahedral 2-fold axis is parallel to the hexagonal axis, but asymmetrically. This OR was also not obtained in our study.

In addition to already known ORs, we have discovered new crystallographic matches between TiB₂ and IQC phase, i.e., hOR2 and hOR3. We have shown that at least one of close-packed planes for TiB₂ exhibits a plane lattice mismatch not greater than ± 2 % with at least one close-packed plane in IQC.

Microstructural analysis showed that after solidification, TiB₂ particles were found in the central part of the primary metastable IQC phase with specific OR, confirming that TiB₂ particles had an effect on the formation of the primary IQC phase. Considering the average IQC volume fraction in BIO1 and BIO1-TiB2 alloys of 18.1 ± 0.6 % and 18.6 ± 0.4 %, respectively, little change was found. The purpose of inoculation is to promote heterogeneous nucleation of the inoculated phase by introducing suitable substrates that serve as nuclei and initiate growth. It is known that inoculation reduces the degree of undercooling needed before the inoculated phase forms; moreover, inoculation prolongs the formation and growth of the inoculated phase and thus the solidification time. The specific content of the phases formed during solidification is influenced primarily by the composition of the alloy, the cooling rate during solidification, and any additional treatments or modifications

applied to the melt. The chemical composition of the alloys studied was almost identical; only the solidification time could be extended in the case of the BIO1-TiB2 alloy by the addition of TiB₂ inoculants. If we take this into account, the primary metastable IQC phase should be larger in the BIO1-TiB2 alloy, as can be seen in the area near the contact surface (Fig. 4a and c), but this is not the case for the central area (Fig. 4b and d), where the IQC phase is larger in the BIO1 alloy. The smaller size of the IQC phase of the BIO1-TiB2 alloy in the central region could be explained by the increased number of nucleation sites in the alloy, resulting in a refinement of the IQC phase. The large difference in the size of the IQC phase between the area near the contact surface and the central area of the BIO1 alloy could also be due to the primary IQC phase starting to grow at the cold mould surface and moving toward the central area. Some of the primary phases might migrate toward the centre with the solidification front, resulting in a larger amount and larger dendritic branches of the IQC phase. On the other hand, the IQC phase in the BIO1-TiB2 alloy grows more homogeneously throughout the entire volume. The higher number of nucleation sites leads to a more uniform distribution and smaller size differences of the IQC phase between the area near the contact surface and the central area.

We believe that the average orientation of the IQC phase in the BIO1-TiB2 alloy changed compared to BIO1 due to the addition of TiB₂ as an inoculant, as quasicrystalline dendrites with a specific orientation
relationship are formed on the TiB₂. TiB₂ particles added to liquid Al have a well-documented crystallographic orientation relationship that provides high coherence, with the planes (0001) in TiB₂ and (111) of αAl being parallel and the directions [1120] in TiB₂ and [110] in αAl phase also being parallel. From Table 2, it can be seen that the plane (0001) and the direction [1120] in TiB₂ are parallel to the 5-fold and 2-fold planes/axes in IQC. This means that the 5-fold and 2-fold planes in IQC should be parallel to the (111) plane and the [110] direction in the αAl phase. Further studies on these alloys will be conducted to confirm this hypothesis.

Compression tests on cylindrical specimens show that the compres-
sive fracture strain increased significantly with the addition of TiB2
particles, while the compressive yield strength and compressive strength
did not change much and remained within the statistical error range
(Graph 1). Considering the small difference in the volume fraction of
IQC phase between the alloys, the results correlate well with the
measured compressive yield strength and compressive strength of the
BIO1 alloy with 390 ± 16 MPa and 616 ± 30 MPa, and the BIO1-TiB2
alloy with 394 ± 15 MPa and 664 ± 23 MPa. On the other hand, it is
evident that the size of IQC phase affects the compressive fracture strain,
which was 20 ± 2 % for BIO1-TiB2 alloy and 12 ± 3 % for BIO1 alloy.
We believe that the shorter dendritic branches of the primary IQC phase
in the central area (Fig. 4) and the more homogeneous distribution of the
IQC phase over the specimen cross-section (Fig. 5) in the BIO1-TiB2
alloy have a significant influence on the reduction of the notch effect
and thus on the higher ductility of this alloy. The compressive yield
strength and compressive strength depend mainly on the grain size of
the Al matrix and the volume fraction and size of the second phase. In
our case, a slightly higher value for BIO1-TiB2 could be due to the
smaller size of the primary IQC phase, while the volume fraction of the
IQC phase and the grain size of the matrix are more or less the same. The
ductility of the alloys depends largely on the homogeneity of the
microstructure. It can be clearly seen that a more homogeneous distri-
bution of the primary IQC phase in the BIO1-TiB2 alloy leads to higher
ductility. In the case of an inhomogeneous microstructure, as in the BIO1
alloy, the weaker regions may be the cause of fractures and reduce
ductility.

**5. Conclusions**

The aim of this study was to investigate and elucidate the nucleation
of the metastable IQC phase in rapidly solidified Al-Mn-Si-based alloys
when their melt is inoculated with TiB₂ particles from the AlTi3B1
master alloy. In summary:

* TiB₂ particles in the AlTi3B1 alloy have a faceted hexagonal platelet morphology with {0001} planes as basal facets and {0110} or {0111} planes as prism facets. This morphology is similar to the theoretical growth morphology calculated from the binding energies of the crystal facets in the Al-based melt [31].

* Metallographic analysis revealed that inoculation of the BIO1 with TiB₂ particles reduces the length of the dendrite arms of the primary IQC phase, contributes to a more homogeneous distribution, and slightly increases the content of the primary metastable IQC phase.

* TiB₂ particles located in the middle of the primary IQC phase confirm that IQC phase nucleates heterogeneously on TiB₂ particle, which served as a substrate. Five orientation relationships were found between TiB₂ and the primary metastable IQC phase. All the orientation relationships found between a crystalline substrate and the metastable IQC phase can be explained by the heterogeneous nucleation of the metastable IQC phase on the facets of the TiB₂, such that the TiB₂ facets are parallel to at least one of the close-packed planes of the metastable IQC phase that have the lowest surface energy (5-fold and 2-fold planes).

* The compression tests show that the addition of TiB₂ particles increases the compressive fracture strain, while compressive strength and the compressive yield strength remains essentially unchanged. An increase in the ductility of the inoculated alloys is associated with a decrease in the notch effect due to the shorter dendritic arms of the primary IQC phase.

CRediT authorship contribution statement

**Blaž Leskovar:** Conceptualization, Investigation, Visualization, Writing – Original Draft. **Sašo Sturm:** Methodology, Writing – Review & Editing. **Kemal Delijić:** Investigation, Validation. **Julian Ledieu:** Validation, Writing – Review & Editing. **Matejka Podlogar:** Investigation. **Sandra Drev:** Investigation. **Iztok Naglic:** Data Curation, Investigation. **Bošjan Markoli:** Conceptualization, Writing – Review & Editing.

Declaration of Competing Interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

Data Availability

Data will be made available on request.

Acknowledgement

The authors thank the Slovenian Research Agency (ARRS) for
financial support under Program No. MR: 38148, P1-0195, P2-0084 and
the International Associated Laboratory – LIA PACS2 funded by INC
CNRS, Université de Lorraine and the Jožef Stefan Institute. Part of this
research was also financially supported by the EIT RawMaterials project
CastQC under project number 21128.

**Appendix A. Supporting information**

Supplementary data associated with this article can be found in the
online version at doi:10.1016/j.jallcom.2023.172195.

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