TiAl composite

Accepted Manuscript Investigation on the crystallographic orientation relationships and interface atomic structures in an in-situ Ti2AlN/TiAl composite

Pei Liu, Dongli Sun, Xiuli Han, Qing Wang PII: DOI: Reference:

S0264-1275(17)30542-7 doi: 10.1016/j.matdes.2017.05.061 JMADE 3085

To appear in:

Materials & Design

Received date: Revised date: Accepted date:

20 February 2017 7 May 2017 22 May 2017

Please cite this article as: Pei Liu, Dongli Sun, Xiuli Han, Qing Wang , Investigation on the crystallographic orientation relationships and interface atomic structures in an in-situ Ti2AlN/TiAl composite, Materials & Design (2017), doi: 10.1016/j.matdes.2017.05.061

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ACCEPTED MANUSCRIPT Investigation on the crystallographic orientation relationships and interface atomic structures in an in-situ Ti2AlN/TiAl composite Pei Liu, Dongli Sun, Xiuli Han, Qing Wang School of Materials Science and Engineering, Harbin Institute of Technology, Harbin15001, People's Republic of china

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Abstract: In this research work, 15 vol.% Ti2AlN particles reinforced TiAl alloy matrix composite was prepared by an in-situ reactive hot processing method using Ti, Al and TiN powders as raw materials. Transmission electron microscope (TEM) coupled with energy-dispersive spectroscopy (EDS) was utilized to investigate the morphology characteristics of Ti2AlN particles. The crystallographic orientation relationships and interface atomic structures between various shaped Ti2AlN particles and TiAl matrix were investigated by the means of selected area electron diffraction (SAED) and high resolution transmission electron microscopy (HRTEM). The results show that four kinds of Ti2AlN particles presented in the composite, which include vast majority of elliptical and polygonal Ti2AlN particles with the sizes of 4μm, few hexagonal Ti2AlN particles with the sizes of 4μm and small plate Ti2AlN particles with the sizes of 0.4μm. The elliptical and polygonal Ti2AlN particles have no fixed crystallographic orientation relationship with TiAl matrix, the elliptical and polygonal Ti2AlN/TiAl interfaces are incoherent interphase boundaries composed by many atomic-scale microfacets. The hexagonal and small plate Ti2AlN particles have a fixed orientation relationship with TiAl matrix. Both the hexagonal and small plate Ti2AlN particles have a ledge and terrace coherent interface with TiAl matrix.

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Key words: Ti2AlN/TiAl composite; reactive hot processing; crystallographic orientation relationship; interface atomic structure

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Corresponding authors at: School of Materials Science and Engineering, Harbin Institute of Technology, 92 West Dazhi Street, Nan Gang District, Harbin15001, PR China E-mail addresses: [email protected] (Dongli Sun), [email protected] (Xiuli Han)

ACCEPTED MANUSCRIPT 1 Introduction

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With the rapid development of the aviation and automobile industry in recent years, there is an increasing requirement in terms of better performance of the high-temperature materials[1]. TiAl alloys have remarkable properties, such as low density, good creep resistance and high strength at elevated temperature, as well as excellent oxidation and corrosion resistance. Hence, it has been regarded as the enormous potential candidate to replace traditional Ni-based superalloys as gas turbine part materials in the field of aircraft engines[2-4]. However, their industrial applications have been limited by several shortcomings, such as intrinsic brittleness and poor oxidation resistance at high temperature. Thus, considerable researches have been devoted to improve their mechanical properties and ductility over the last decades, such as heat treatment[5-6], thermal-mechanical processing[7], alloying[8-9] and composite technology[10-20]. Bolz et al.[6] investigated the effect of two-step heat treatment on the microstructure and mechanical properties of a TiAl alloy, and reported that the optimum condition is heating treating the material at 1543K followed by air cooling and annealing at 1073K. Shu et al.[8] studied the effects of Fe, Co, Ni elements on the ductility of TiAl alloy, finding that the tested average fracture strain of TiAl alloy increases from 17.3% to 19.1% and 18% with the addition of 3 at.% Fe and Co, respectively. Sun et al.[14] synthesized Ti2AlC/TiAl composites utilizing the method of hot press consolidation followed by isothermal forging, and reported that the as-forged composite shows better ductility without sacrificing strength, with elongation (δ) increasing 12.10% to 30.87%. Compared with other methods, composite technology has attracted much attention due to its good combination of the property of matrix and reinforcement. In the researches of TiAl matrix composites, ceramic particles, like TiB2[10], Ti5Si3[11], Al2O3[12], TiC[13], Ti2AlC[14-16] and Ti2AlN[17-20] have been identified as compatible and thermochemically stable reinforcing phases for the TiAl matrix. Among these ceramic particles, ternary compounds, like Ti2AlC and Ti2AlN, are fascinating members of the family of MAX phases with the general formula of Mn+1AXn(n=1,2,3). These ternary compounds are a series of new type of machinable ceramic materials, and possess the extraordinary properties of both metallic and ceramic characteristics. The distinctive combination of these properties stems from co-existence of the strong covalent-ionic M-X bonds and the weak metallic M-A bonds inside the layered hexagonal structures of MAX materials[21-22]. Compared with Ti2AlC, Ti2AlN displays superior mechanical properties [23], so it is a competitive and potential reinforcement for TiAl matrix alloy. The processing techniques to fabricate TiAl matrix composites could be divided into in-situ and ex-situ process according to the addition modes of reinforcing particles. Compared with ex-situ process, in-situ process possesses some merits, e.g., perfect reinforcement-matrix interface, much lower cost of production [15, 24-25]. Many processing techniques have been developed to fabricate in-situ TiAl matrix composites, such as mechanical alloying[13, 26], vacuum arc-melting[20], combustion synthesis[16], and reactive hot processing[14, 17-19]. Reactive hot pressing process can adjust the composition, fraction and distribution of the reinforcing particles, hence it has become an attractive method to fabricate TiAl matrix composite. In recent years, many studies have been performed on Ti2AlN/TiAl composite with the objective of providing the fundamental knowledge essential to the development of this new composite. Our previous researches[17-19] have investigated the hardness, elastic modulus, high-temperature compressive strength and wear behavior of Ti2AlN/TiAl composite, and the

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results showed that the mechanical properties and wear resistance of reactive hot pressing synthesized Ti2AlN/TiAl composite are improved obviously compared with TiAl matrix. Liu et al.[20,27] investigated the tensile properties of Ti2AlN/TiAl composites at room and elevated temperatures, finding that 4 vol.% Ti2AlN/TiAl composite exhibits the fracture strength of 670 MPa, fracture strain of 0.39% at room temperature and tensile strength of 645.3 MPa at 1073K, which are 71%, 14.7% and 70.5% higher than those of TiAl alloy, respectively. Kakitsuji et al.[28] investigated the bending strength of Ti2AlN/TiAl composites, and reported that the bending strength and fracture toughness of these composites were larger than that of TiAl alloy, and the ductility kept the same level with that of TiAl alloy. Interface is an extremely important component of metal matrix composites, the microstructure of an interface is a key factor that influence the load transfer, strengthening mechanism and fracture process, determining if the matrix and reinforcement particles can give full play to their roles and form the optimal comprehensive performance[29-30]. In addition, the synthesis of reinforcement particles in an in-situ composite is a phase transformation process, in which the interface can act as a “barrier” or “channel” for phase transformation[31-32], thus the nature and behavior of interface in an in-situ composite play an important role in the nucleation and growth of reinforcement particles. Undoubtedly, understanding of the crystallographic orientation relationships and interface atomic structures between TiAl matrix and Ti2AlN particles is of great significance in guiding the preparation and application of high-performance Ti2AlN/TiAl composites. There have been three literatures[33-35] reported the crystallographic orientation relationship between Ti2AlN and TiAl. Tian et al.[33] investigated the precipitation behavior of nitrides in L10-ordered TiAl and reported a — crystallographic orientation relationship between Ti2AlN precipitates and TiAl matrix, i.e.[112 — 0]Ti2AlN//[101]TiAl, (0001)Ti2AlN//(111)TiAl. Cui et al.[34] synthesized Ti2AlN ceramic by spark plasma sintering at 1473K, and found the same crystallographic orientation relationship between Ti2AlN nanowhiskers and TiAl. Chiu et al. [35] investigated the atomic configurations of nitride phase at the interface of aluminum nitride and titanium foil after annealing at 1673K, and reported another — — crystallographic orientation relationship between Ti2AlN and TiAl, i.e.[1120]Ti2AlN//[110]TiAl, (110 — —— 3)Ti2AlN//(111)TiAl. It can be inferred from the above three investigations that the crystallographic orientation relationship between Ti2AlN and TiAl varies with the experimental conditions. What is more, to the authors’ knowledge, there have been no reports about the crystallographic orientation relationships and interface atomic structures between Ti2AlN particles and TiAl matrix in Ti2AlN/TiAl composite. Based on the aforementioned investigations, the properties of Ti2AlN/TiAl composite are very promising. If it is to be selected as a new type of high temperature structural material and applied in industries, the substantial problem and great challenge are the further optimization of comprehensive mechanical performance. Their optimization requires fundamental investigations of interface microstructural characteristics and their effects on the formation of Ti2AlN particles during the preparation, and of the relationship between interface microstructure and mechanical properties. The aim of this work is to understand comprehensively the crystallographic orientation relationship and interface atomic structures between Ti2AlN particles and TiAl matrix in Ti2AlN/TiAl composite and provide the theoretical basis for the material design, performance optimization and engineering application of Ti2AlN/TiAl composites. In the present paper, 15 vol.% Ti2AlN/TiAl composite was fabricated by an in-situ reaction hot-pressing sintering method. Transmission electron microscope (TEM) coupled with energy-dispersive spectroscopy (EDS) was utilized to investigate the

ACCEPTED MANUSCRIPT morphology characteristics of Ti2AlN particles. The crystallographic orientation relationships and interface atomic structures between various shaped Ti2AlN particles and TiAl matrix are investigated by the means of selected area electron diffraction (SAED) and High Resolution Transmission Electron Microscopy (HRTEM).

2 Experimental procedures

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The Ti2AlN/TiAl composite used in this study for investigation was manufactured by powder metallurgy route. Ti powders (99.4% purity, mean particle size was less than 25µm), Al powders (99.4% purity, mean particle size was less than 5µm) and TiN powders (99.4% purity, mean particle size was less than 3µm) were used as raw materials. The powders with a deigned composition (Ti:Al:TiN=11.01:9.88:1 in molar ratio) were mixed and then were milled for 12h with addition of ethanol using Al2O3 grinding media, the ball-to powder mass ratio is 5:1 and revolving speed is 120r/min. After milling, powders were dried in a vacuum oven. As confirmed by numerous literatures[17-19,34], the best temperature for preparing Ti2AlN ceramic or Ti2AlN/TiAl composite using the powder metallurgy process is between 1573-1673K, because the high purity Ti2AlN particles cannot be acquired when the temperature is below 1573K and the Ti2AlN can decomposed into Ti4AlN3 through the sublimation of Al when the temperature is above 1673K. In addition, the complicated synthesis process of Ti2AlN includes three major steps: firstly, Al powders melt and react with Ti powder to form TiAl3 at 973-1173K; secondly, Ti-Al intermetallics react with the remaining Ti cores to form Ti3Al and TiAl phases at 1173-1473K; finally, the TiN powers react with TiAl phase and result in a plenty of Ti2AlN phase at 1473-1673K. Thus, in order to acquire the Ti2AlN/TiAl composite with the designed composition, the sintering process in our experiment is elaborately designed as follows: the mixed powders were put into BN-coated graphite die and were heated to 903K at the heat rate of 10K/min with an axial pressure of 10MPa. After holding 1h at 903K, the powders were heated to 1623K at the heat rate of 6K/min with no pressure. When the temperature arrived 1623K, an axial pressure of 25MPa was applied and held for 1h, then releasing the pressure and cooling in the furnace to ambient temperature, the target material was achieved. Microstructural characterization was performed in the hot-pressed states. The samples for TEM observations were machined into 0.5mm by wire-electrode cutting and ground to 50μm by Mechanical thinning, and then were cut into 3mm diameter foils. After that, foils were prepared by argon ion milling using Gatan 691 Precision Ion Polishing System. Transmission electron microscopy (TEM) investigations were performed on Talos F200x field emission transmission electron microscopy coupled with EDS at an accelerating voltage of 200 kV.

3 Experimental Results 3.1 Morphology of Ti2AlN Particles Fig. 1 shows the TEM morphology, EDS spectra and SAED pattern of Ti2AlN particles in the in-situ Ti2AlN/TiAl composite fabricated by the reactive hot-processing method. It can be seen from Fig.1 that the Ti2AlN particle have four morphology categories: elliptical Ti2AlN, polygonal Ti2AlN, hexagonal Ti2AlN and small plate Ti2AlN. Most of the Ti2AlN particles are elliptical and polygonal, the particle size is about 4μm.There are also small quantities of hexagonal Ti2AlN particles and

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small plate Ti2AlN particles, the hexagonal Ti2AlN particle size is about 4μm, and the small plate Ti2AlN particle size is about 0.4μm. It is worth pointing out that the small plate Ti2AlN particles are more likely to attach itself to the polygonal Ti2AlN particles. The reason why the small plate Ti2AlN particles are more likely to attach itself to the polygonal Ti2AlN particles will be discussed in a later section 4.2.2.

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Fig.1 TEM morphology, EDS spectra and SAED pattern of Ti2AlN particles in the in-situ Ti2AlN/TiAl composite fabricated by the reactive hot-processing method (a) Elliptical Ti2AlN；(b) Polygonal Ti2AlN；(c) Hexagonal Ti2AlN；(d) Small plate Ti2AlN

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3.2 The interfacial microstructure between Ti2AlN and TiAl 3.2.1 Elliptical Ti2AlN/TiAl interface Fig.2(a) shows a typical TEM image of the elliptical Ti2AlN particle in-situ synthesized in this experiment, the interface between Ti2AlN particle and TiAl matrix is clean and there are no interface reactants. Fig.2(b) shows the selected area electron diffraction (SAED) pattern from the elliptical Ti2AlN particle and TiAl matrix. It is worth pointing that there are some small extra spots in addition to the main diffraction spots of Ti2AlN and TiAl in Fig.2(b), these small extra spots mainly came from double diffraction. Fig.2(c) is corresponding indexed pattern without small extra spots, the blue dots represent the reflections from the Ti2AlN, and the red dots represent the — reflections from TiAl. As indexed in Fig.2(c), the incident beam is parallel to [2111]Ti2AlN and [13 — — 2]TiAl, and the (1013) plane of Ti2AlN particle is parallel to the (112) plane of TiAl matrix. Thus,

ACCEPTED MANUSCRIPT their orientation relationship, denoted as OR I in this paper, is determined as follows: — — — [2111]Ti2AlN // [132]TiAl, (1013)Ti2AlN // (112)TiAl —

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A HRTEM image along the [2111]Ti2AlN // [132]TiAl is shown in Fig.2(d), where the Ti2AlN phase is at the right side of the image, and the TiAl is at the left side, with their Fast Fourier Transform(FFT) patterns inserted. As seen in the FFT of Ti2AlN, there are four satellite diffraction spots in each diffraction spot of Ti2AlN, it is just these satellite diffraction spots that make the periodic moiré fringes appear in the HRTEM image of Ti2AlN. The Inverse Fast Fourier Transform (IFFT) image of black square area in Fig.2(d) is showed in Fig.2(e), the blue dots represent the atoms of Ti2AlN and the red dots represent the atoms of TiAl. It can be seen clearly from Fig.2(e) — that the smooth curve of Ti2AlN surface is not atomically flat, but comprised of many (1013) and (1 — — 103) microfacets. The (1013) microfacet of Ti2AlN particle is parallel to the (112) microfacet of TiAl matrix, but there is no common atomic plane or the periodic mismatch dislocation between — them, so the interface between (1013)Ti2AlN microfacet and (112)TiAl microfacet is incoherent. The (1 — — — 103) Ti2AlN microfacet is not parallel to the (111) TiAl microfacet, there is a rotation angle of 10° between these two microfacets, which means that this interface is incoherent. Therefore, we can conclude that the smooth curved interface between elliptical Ti2AlN particle and TiAl matrix is an incoherent interphase boundary composed by many atomic-scale microfacets.

Fig.2 The crystallographic orientation relationship and interface atomic structure between elliptical Ti2AlN particle and TiAl matrix (a) TEM image of elliptical Ti2AlN particle; (b) SAED pattern from the elliptical Ti2AlN/TiAl interface showing the OR I; (c) Indexing of the SAED pattern in Fig.2(b); (d) HRTEM image of the Ti2AlN/TiAl interface along the black

ACCEPTED MANUSCRIPT square area in Fig.2(a); (e) IFFT image of the black square area in Fig.2(d)

Fig.3(a) shows another TEM image of the elliptical Ti2AlN particle in-situ synthesized in this experiment. A second type of orientation relationship, OR II, is observed, as shown in Fig.3(b). The — — Ti2AlN is along the [2111] zone axis while the TiAl is along the [101] zone axis, as indexed in Fig.3(c). Thus the following OR II results: — — — [2111]Ti2AlN // [101]TiAl, (1013)Ti2AlN // (111)TiAl —

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A HRTEM image along the [2111]Ti2AlN // [101]TiAl is shown in Fig.3(d), where the Ti2AlN phase is located on the top of the TiAl, with their FFT patterns inserted. Along their interface, the left side appears periodic moiré fringes, as marked in the image by the white square. Fig.3(e) is the IFFT image of the white square area in Fig.3(d), and it can be seen clearly that the spacing of moiré fringes is about 0.83nm. It is well known that the formation of moiré fringes between two superimposed crystal lattices of different lattice space can be explained by double diffraction. In general the spacing of moiré fringes is given by the expression[36]: d  d  2d1d 2 cos  2 1

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where d1 and d2 are the relevant lattice plane spacing, φ is the angle between them. The observed periodic moiré fringes can be explained quantitatively by reference to Fig.3(b). — — — Based on the diffraction pattern shown in Fig.3(b), the interaction of (1103)Ti2AlN and (111)TiAl is — — responsible for the double diffraction. The interplanar spacing of (1103)Ti2AlN is 0.225nm, the — interplanar spacing of (111)TiAl is 0.232nm, the angle between them is 15°, thus the moiré spacing can be calculated as D=0.87nm according to the formula(1), which agrees well with the measured value of 0.83nm from HRTEM image in Fig.3(e). As shown in Fig.3(d), the right side of the interface has no moiré fringes. Fig.3(f) is the IFFT image of black square area in Fig.3(d), the blue dots represent the atoms of Ti2AlN and the red dots represent the atoms of TiAl. It can be seen clearly from Fig.3(f) that the smooth curve of Ti2AlN — — — — surface is also not atomically flat, but comprised of many (1013) and (1103) microfacets. The (1013) microfacet of Ti2AlN particle is parallel to the (111) microfacet of TiAl matrix, but there are no common atomic plane or the periodic mismatch dislocations between them, so the interface — — — between (1013)Ti2AlN microfacet and (111)TiAl microfacet is incoherent. The (1103) Ti2AlN microfacet — — is not parallel to the (111) TiAl microfacet, there is a rotation angle of 15° between these two microfacets. Therefore, the smooth curved interface between this elliptical Ti2AlN and TiAl is also an incoherent interphase boundary composed by atomic-scale microfacets.

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Fig.3 The crystallographic orientation relationship and interfacial atomic structure between elliptical Ti2AlN particle and TiAl matrix (a) TEM image of elliptical Ti2AlN particle; (b) SAED pattern from the elliptical Ti2AlN/TiAl interface showing the OR II; (c) Indexing of the SAED pattern in Fig.3(b); (d) HRTEM image of the Ti2AlN/TiAl interface along the black square area in Fig.3(a); (e) The IFFT image of the white square area in Fig.3(d); (f) The IFFT image of the black square area in Fig.3(d)

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3.2.2 The polygonal Ti2AlN/TiAl interface

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Fig.4(a) shows a typical TEM image of the polygonal Ti2AlN particle in-situ synthesized in this experiment. Compared with the elliptical Ti2AlN particle containing only curved surface, the polygonal Ti2AlN particle contains both smooth curved surface and flat surface. Fig.4(b) shows the selected area electron diffraction (SAED) pattern from the polygonal Ti2AlN particle and TiAl — — matrix. As indexed in Fig.4(c), the incident beam is parallel to [2111]Ti2AlN and [101]TiAl , and the (10 — 1 3) plane of Ti2AlN particle is parallel to the (111) plane of TiAl matrix. Their orientation relationship is same as OR II, namely: — — — [2111]Ti2AlN // [101]TiAl, (1013)Ti2AlN // (111)TiAl A HRTEM image along the flat interface is shown in Fig.4(d), where the Ti2AlN phase is at the bottom side of the image, and the TiAl is at the top side, with their Fast Fourier Transform(FFT) patterns inserted. The IFFT image of white square area in Fig.4(d) is showed in Fig.4(e), the blue dots represent the atoms of Ti2AlN and the red dots represent the atoms of TiAl. It can be seen clearly from Fig.4(e) that the flat surface of Ti2AlN is not atomically flat, but comprised of many — — — (1013), (1103) and (0110) microfacets. The interphase interface between Ti2AlN particle and TiAl — — — 15° — — — 5° matrix is composed by (1013)Ti2AlN//(111)TiAl, (1103) Ti2AlN ~ (111) TiAl and (0110) Ti2AlN ~ (020)

ACCEPTED MANUSCRIPT , which is agreed with the indexed diffraction pattern in Fig.4(c). There are no common atomic planes or periodic mismatch dislocations between all the microfacets of polygonal Ti2AlN particle and TiAl matrix, so the smooth flat interface between polygonal Ti2AlN particle and TiAl matrix is an incoherent interphase boundary composed by many atomic-scale microfacets.

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Fig.4 The crystallographic orientation relationship and interfacial atomic structure between polygonal Ti2AlN particle and TiAl matrix (a) TEM image of polygonal Ti2AlN particle; (b) SAED pattern from the polygonal Ti2AlN/TiAl interface showing the OR II; (c) Indexing of the SAED pattern in Fig.4(b); (d) HRTEM image of the Ti2AlN/TiAl interface along the white square area in Fig.4(a); (e) IFFT image of the white square area in Fig.4(d)

Fig.5(a) shows another TEM image of the polygonal Ti2AlN particle in-situ synthesized in this experiment. A third type of OR, OR III, is observed, as shown in Fig.5(b). The Ti2AlN is along the — — [0331] zone axis while the TiAl is along the [110] zone axis, as indexed in Fig.5(c). Thus the following OR III results: — — — [0331]Ti2AlN // [110]TiAl, (1013)Ti2AlN // (110)TiAl A HRTEM image along the flat interface is shown in Fig.5(d), where the Ti2AlN phase is at the bottom side of the image, and the TiAl is at the top side, with their Fast Fourier Transform(FFT) patterns inserted. The IFFT image of black square area in Fig.5(d) is showed in Fig.5(e), the blue dots represent the atoms of Ti2AlN and the red dots represent the atoms of TiAl. It can be seen clearly from Fig.5(e) that the flat surface of Ti2AlN is not atomically flat, but comprised of many — — — (1013 and (1103) microfacets. The interphase interface between Ti2AlN particle and TiAl matrix is — — — 10° composed by (1013)Ti2AlN//(110)TiAl and (1103) Ti2AlN ~ (001) TiAl, which is agreed with the indexed

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diffraction pattern in Fig.5(c). There are no common atomic planes or the periodic mismatch dislocations between all the microfacets of polygonal Ti2AlN particle and TiAl matrix, so the smooth flat interface between this polygonal Ti2AlN particle and TiAl matrix is also an incoherent interphase boundary composed by many atomic-scale microfacets, which is same as the results shown in Fig.4.

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Fig.5 The crystallographic orientation relationship and interface atomic structure between polygonal Ti2AlN particle and TiAl matrix (a) TEM image of polygonal Ti2AlN particle; (b) SAED pattern from the polygonal Ti2AlN/TiAl interface showing the OR III; (c) Indexing of the SAED pattern in Fig.5(b); (d) HRTEM image of the Ti2AlN/TiAl interface along the white area in Fig.5(a); (e) IFFT image of the white square area in Fig.5(d)

3.2.3 Hexagonal Ti2AlN/TiAl interface Fig.6(a) shows a TEM image of the hexagonal Ti2AlN particle in-situ synthesized in this experiment. A nearly perfect orientation relationship, OR IV, is observed, as shown in Fig.6(b). This orientation relationship is consistent with previous work on the nitride precipitates in L10-ordered — TiAl[33] and formation of Ti2AlN nanowhiskers in Ti2AlN ceramics[34]. The Ti2AlN is along the [112 — 0] zone axis while the TiAl is along the [101] zone axis, as indexed in Fig.6(c). Thus the following ORIV results: — — — — — — — [1120]Ti2AlN // [101]TiAl, (0006)Ti2AlN // (111)TiAl, (1104)Ti2AlN // (020)TiAl, (1102)Ti2AlN // (111)TiAl A HRTEM image along the interface between the top edge of hexagonal Ti2AlN and TiAl is

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shown in Fig.6(d), where the Ti2AlN phase is at the lower right side of the image, and the TiAl is at the upper left, with their Fast Fourier Transform(FFT) patterns inserted. It can be seen that the boundaries of the Ti2AlN are (0001) surfaces with ledges. An IFFT of the structure ledges (the white square area in Fig.6d) is shown in Fig.6(e). The atomic stacking sequence of Ti2AlN is visible from the lower side of Fig.6(e), and can be described as the sequence of BABABA along the [0001] direction, where the underlined letters refer to Al layers and the rest to Ti layers. It can also be seen — clearly from Fig.6(e) that the terrace is a (0001)Ti2AlN//(111)TiAl interface, the ledge is a (1 104)Ti2AlN//(020)TiAl interface. Both the terrace and ledge interfaces display good atomic matching, which means that these interfaces are coherent. The ledges on the interface indicate that the growth of hexagonal Ti2AlN particle could be described by a ledge growth and the lateral movement of the monoatomic layers at the interface.

Fig.6 The crystallographic orientation relationship and interface atomic structure between hexagonal Ti2AlN particle and TiAl matrix (a) TEM image of hexagonal Ti2AlN particle; (b) SAED pattern from the hexagonal Ti2AlN/TiAl interface showing the OR IV; (c) Indexing of the SAED pattern in Fig.6(b); (d) HRTEM image of the Ti2AlN/TiAl interface along the white square area in Fig.6(a); (e) IFFT image of the white square area in Fig.6(d)

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3.2.4 Small plate Ti2AlN/ TiAl interface

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A type of small plate Ti2AlN particle is also observed, although less frequently, as shown in Fig.7(a). The small plate Ti2AlN particle is more likely to attach itself to the larger polygonal Ti2AlN particle, which implies that the growth mechanism of small plate Ti2AlN might be different from that of polygonal Ti2AlN particle. The diffraction patterns of polygonal Ti2AlN particle and small plate Ti2AlN particle are respectively shown in Fig.7(b) and Fig.7(c), the indexed results — show that the polygonal Ti2AlN particle is along the [2111] zone axis and the small plate Ti2AlN — particle is along the [1120] zone axis. A magnified image of the small plate Ti2AlN is shown in Fig.7(d). Fig.7(e) shows the selected area electron diffraction (SAED) pattern from the small plate — Ti2AlN particle and TiAl matrix. As indexed in Fig.7(f), the incident beam is parallel to [1120]Ti2AlN — and [101]TiAl. Their OR is same as the nearly perfect OR IV shown in Fig.6(b), namely: — — — — — — — [1120]Ti2AlN // [101]TiAl, (0006)Ti2AlN // (111)TiAl, (1104)Ti2AlN // (020)TiAl, (1102)Ti2AlN // (111)TiAl

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A HRTEM image along the interface of small plate Ti2AlN and TiAl is shown in Fig.7(g), where the Ti2AlN phase is at the right side of the image, and the TiAl is at the left, with their Fast Fourier Transform(FFT) patterns inserted. It can be seen that the boundaries of the Ti2AlN are (0001) surfaces with ledges. An IFFT of the structure ledges (the black square area in Fig.7g) is shown in Fig.7(h) and the interface displays good atomic matching, suggesting that small plate Ti2AlN particle has a good coherency with TiAl matrix. Since both the orientation relationship and interface atomic structure shown in Fig.7 are same as those shown in Fig.6, the growth of small plate Ti2AlN particle could also be described by a ledge growth and the lateral movement of the monoatomic layers at the interface.

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Fig.7 The crystallographic orientation relationship and interface atomic structure between small plate Ti2AlN particle and TiAl matrix (a) TEM image of small plate and polygonal Ti2AlN particle; (b) Diffraction pattern from the polygonal Ti2AlN particle; (c) Diffraction pattern from the small plate Ti2AlN particle; (d) A magnified image of the small plate Ti2AlN particle; (e) SAED pattern from the small plate Ti2AlN/TiAl interface showing the OR IV; (f) Indexing of the SAED pattern in Fig.7(e); (g) HRTEM image of the Ti2AlN/TiAl interface along the white square area in Fig.7(d);(h) IFFT image of the black square area in Fig.7(g)

4 Discussions 4.1 Analysis of observed ORs between Ti2AlN and TiAl

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u  h  v   G 1  k       w l 

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As shown in section 3, four types of crystallographic orientation relationships have been observed between Ti2AlN particles and TiAl matrix in our experiment. Since a crystallographic orientation relationship can be described in many formats according to the symmetry principle in the crystallography, comparison between different types of crystallographic orientation relationship should be done. Transformation matrix is an effect method to distinguish different variants of the crystallographic orientation relationships. The center idea of matrix method is to calculate the transformation matrix B and A of every crystallographic orientation relationship. If the absolute values of the nine elements in the transformation matrix are same, in spite of the orders and positions are not same, then it can be concluded that these two crystallographic orientation relationships belong to the same type. According to literature[37], crystal plane (hkl) can be expressed in the reciprocal space by a * reciprocal vector Ghkl  ha1*  ka2*  la3* . The normal direction of crystal plane (hkl) is the positive vector and can be expressed by I uvw  ua1  va2  wa3 . The relationship * * * * between Ghkl  ha1  ka2  la3 and I uvw  ua1  va2  wa3 is as follows:

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h  u   k   G v      l   w

For any crystal system,

ab cos  b2 bc cos 

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 a2  G   ab cos   ac cos  

ac cos    bc cos   c 2 

(2)

(3)

(4)

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Where a, b, c, α, β, γ are the lattice parameter. In the general crystallographic study, the orientation relationship between two phases is always expressed in the form of [u'2v'2w'2]//[u2v2w2] and (h'1k'1l'1)// (h1k1l1). According to [u'2v'2w'2]//[u2v2w2] and formula (2), we can obtain the second crystal plane parallel relationship (h'2k'2l'2)// (h2k2l2). According to (h'1k'1l'1)// (h1k1l1) and formula (3), we can obtain the second crystal orientation parallel relationship[u'1v'1w'1]//[u2v2w2]. Then the (h'3k'3l'3) can be obtained by [u'1v'1w'1]×[u'2v'2w'2], the(h3k3l3) can be obtained by [u1v1w1]×[u2v2w2]. Thus, three crystal plane parallel relationships between the two phases are obtained:  (h1k1l1 ) / /(h1' k1'l1' )  ' ' ' (h2 k2l2 ) / /(h2 k2l2 )  (h k l ) / /(h ' k ' l ' ) 3 3 3  3 33

(5)

The above orientation relationships can be expressed in the form of matrix: u '  u   '   v   B v   w'   w  

or

 h'  h   '   k   A k  l '  l   

(6)

Where B is transformation matrix, A is the transposed inverse matrix of B, i.e. A=(B-1)T, the transformation matrix B can be expressed in the following form:

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 h1'  B   h2'  h3' 

k1' k2' k3'

l1'   l2'  l3' 

1

 d1  '  d1  0   0 

0 d2 d 2' 0

 0   h1  0   h2  h  3 d3   d3' 

k1 k2 k3

l1  l2  l3 

(7)

0   8.964036 4.482018    4.482018 8.964036 0   0 0 185.2321 0 0  16.04   0 16.04 0   0 0 16.5649

(8)

(9)

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GTiAl

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GTi2 AlN

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Where d1, d2, d3 are the interplanar spacing of (h1k1l1), (h2k2l2) and (h3k3l3); d'1, d'2, d'3 are the interplanar spacing of (h'1k'1l'1), (h'2k'2l'2) and (h'3k'3l'3). In our experiment, Ti2AlN is hexagonal crystal and the lattice parameters are: a=b=2.994 Å, c=13.610 Å, α=β=90°, γ=120°, TiAl is tetragonal crystal and the lattice parameters are: a=b=4.005 Å, c=4.070 Å, α=β=γ=90°. Putting these parameters into the formula (4), we can obtain the transformation matrix of Ti2AlN and TiAl from crystallographic orientation to crystallographic plane:

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Table 1 shows four types of crystallographic orientation relationships between Ti2AlN particles and TiAl matrix observed in our experiment. According to the formula (2) ~ (9), the parallel planes and the transformation matrix of different crystallographic orientation relationships can be calculated. As shown in Table 2, the calculated results indicate that these four orientation relationships are different, nor they are equal with the rule of symmetry. Therefore, it can be concluded that the elliptical and polygonal Ti2AlN particles have no fixed or preferential crystallographic orientation relationship with TiAl matrix, while the hexagonal and small plate — Ti2AlN particles were found to have a fixed orientation relationship with TiAl matrix, i.e. [112 — 0]Ti2AlN // [101]TiAl, (0006)Ti2AlN // (111)TiAl.

ORs

[uvw]Ti2AlN//[uvw]TiAl —

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Table1 The orientation relationships between in-situ formed Ti2AlN particle and TiAl matrix —

[2 111]Ti2AlN//[132 ]TiAl —

2

—

Morphology of Ti2AlN Particle

(101 3) Ti2AlN//(112)TiAl

0.267

Elliptical Ti2AlN

—

(101 3) Ti2AlN//(111)TiAl

0.032

Elliptical Ti2AlN, Polygonal Ti2AlN

[033 1]Ti2AlN//[11 0]TiAl

(101 3) Ti2AlN//(110)TiAl

0.201

Polygonal Ti2AlN

[112 0]Ti2AlN//[1 01]TiAl

(0006) Ti2AlN//(111)TiAl

0.024

Hexagonal Ti2AlN, Small plater Ti2AlN

—

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Misfit data

—

[2 111]Ti2AlN// [1 01]TiAl —

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[hkl]Ti2AlN//[hkl]TiAl

—

—

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Table 2 The parallel planes and the transformation matrixs between Ti2AlN particle and TiAl matrix ORs

(hkl)Ti2AlN//(hkl)TiAl

Transformation matrix B

Transformation matrix A

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(2 1 1 1 4 ) Ti2AlN//(132 )TiAl —

—

(13 23) Ti2AlN//(4 21)TiAl —

2

—

(2 1 1 1 4 ) Ti2AlN //(1 01)TiAl —

—

(13 23) Ti2AlN//(12 1)TiAl —

3

—

(0001) Ti2AlN//(111)TiAl

0.75

0.24



0.45

  0.55

0.15

0.27 

0.18 

1.94

2.75 



A   0.15 0.58

 0.13

0.13

  3.28

0.44 

0.71 

0.39

0.75 



0.27 

0.22

0.25 

 0.001 0.53 0.53  A   0.53 0.53 0.001  1.96 1.96 1.96 

 0.63 0.63 1.27  0.63 0.63    0.17 0.17 0.17 

 0.001 0.53 0.53  A   0.53 0.53 0.001  1.96 1.96 1.96 



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(112 0) Ti2AlN//(1 01)TiAl

 0.55

A   0.69

0.37

0.43 1.17 

B   1.27

—

—

0.16

  0.05

(011 5) Ti2AlN//(11 0)TiAl (1 104) Ti2AlN//(001)TiAl

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0.79

B   1.2

—

—

  0.06

0.84

1.187

(101 3) Ti2AlN//(110)TiAl —

1.22

 0.17 0.94 1.20  B   0.23 1.48 0.23    0.29 0.03 0.06 

(101 3) Ti2AlN//(111)TiAl — — ——

 0.56

B   0.96

—

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—

(101 3) Ti2AlN//(112)TiAl

—

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(11 00) Ti2AlN//(12 1)TiAl

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4.2 Analysis of observed interface atomic structure between Ti2AlN and TiAl 4.2.1 The incoherent interface

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As shown in section 3, most of the Ti2AlN particles in our experiment are elliptical and polygonal. Both the elliptical Ti2AlN/TiAl curved interfaces and polygonal Ti2AlN/TiAl flat interfaces are incoherent interphase boundaries composed by many atomic-scale microfacets, displaying poor atomic matching across the interface. In order to determine whether the energies of incoherent interfaces observed in our experiment are low or not, we evaluate the surface energies of various planes using density functional theory (DFT) which was implemented by Cambridge Sequential Total Energy Package (CASTEP). The plane-wave ultrasoft pseudopotentials were employed to present the interactions of electrons with ion cores. The generalized gradient approximation by Perdew and Wang (GGA-PW91) was chosen to represent the exchange-correlation energy. The maximum cutoff energy of the plane wave in the reciprocal space is 360eV, and the k-point sampling grids obtained by using Monkhorst-Pack method were set to 10×10×2 for the bulk and all the slabs, respectively. The Broyden-Fletcher-Goldfarb-Shannon(BFGS) algorithm was applied to relax the models to reach the optimized structures. The energy change during the optimization finally converged to less than 10-5eV/atom, and the force acting on the atoms was less than 0.03eV/Å. The surface energy(Esurf) was calculated using the following expression[38]: Esurf  ( Eslab  i Ni i ) 2 A

(10)

Where Eslab is the total energy of a slab, Ni is the atomic numbers of element i, μi is the chemical potential of element i, A is the surface area. Surface was created by cleaving from the optimized bulk structure and was relax by optimizing all the atoms. The derived surface energies are listed in Table 3. It can be seen that the surface energies of — (1010)-TiAl terminated and (1013) Ti2AlN surfaces are relatively lower. These features indicate that — — the atomic faceting along the low-energy{1013} and {1010}Ti2AlN surfaces play an important role —

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Table 3 Surface energy for various Ti2AlN planes Surface energy(J/m2)

(0001)-Al terminated

1.835

(0001)-Ti terminated

3.726

(0001)-N terminated

2.294

—

(1010)-TiAl terminated

1.267

—

(1010)-N terminated

2.766

—

2.195

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(1120)

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Ti2AlN planes

—

(1013)

2.017

[18]

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In our previous research , the synthesis process of in-situ Ti2AlN/TiAl composites in our experiment has been confirmed and the Ti2AlN particles are mainly formed by the reaction between TiN and TiAl at 1473-1623K. Because the sintering temperature (1623K) is lower than the melting point of TiAl (1773K) and TiN (3223K), the main formation mechanism of Ti2AlN particles is solid-state reactive diffusion. Thus, the elliptical and polygonal Ti2AlN particles could be considered to form through a solid-state reactive diffusion mechanism. There have been some literatures[39-42] reported the effect of incoherent interphase interfaces on the diffusional solid-solid phase transformations and on the mechanical properties of composites. Li et al.[39] acquired a Zr-33.9 at.%N alloy by annealing a Zr foil under an atmosphere of high-purity N at 1473K for 2h, finding that the Zr and ZrN phases display poor atomic matching across the {111} ZrN interface, atomic faceting in the ZrN phase is the dominant factor in determining the {111}ZrN interface plane. Massalski et al.[40] have claimed that incoherent interfaces can be involved in the initiation and propagation of diffusional solid-solid phase transformations. Faceting of the new phase along an incoherent boundary segment during the nucleation stage can lower the nucleation barrier and facilitate the atomic diffusion during the growth. Bartolomé et al.[41] acquired ZrO2-Nb ceramic matrix composites by hot-pressing at 1773K and investigated the influence of ceramic-metal interface on crack growth resistance, and the results showed that the (110)ZrO2-(110)Nb interface is faceted along the (110) planes of Nb grains and atomically incoherent. This interface has a moderate interface strength and are more beneficial for toughening the composite than the interface with too high strength or too low strength, because if the interface is very strong, a high degree of constraint can lead to a triaxial state of stress, resulting in a brittle failure of the ductile reinforcement. Consequently, there would not be a significant increase in the composite toughness under these circumstances. Conversely, for very low reinforcement-matrix interface strength, the reinforcement would readily debond and there would be no crack surface bridging action, again resulting in a limited toughness improvement. Karnesky et al.[42] investigated the strengthening mechanism in aluminum containing incoherent Al2O3 particles, and the results showed that the incoherent ceramic particles distributed within a coarse-grained metallic matrix provide high strength at ambient and elevated temperatures, as they impede dislocation glide and climb. Based on the aforementioned investigations, it can be inferred that the incoherent interphase boundaries composed by many atomic-scale microfacets may facilitate the nucleation and growth of

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Few hexagonal Ti2AlN particles and small plate Ti2AlN particles are also observed in our experiment. Both the hexagonal Ti2AlN particles and small plate Ti2AlN particles have a ledge and terrace coherent interface with TiAl matrix, displaying good atomic matching across the interface. The Ti2AlN/TiAl composite in this experiment is fabricated using the Ti powders, Al powders and TiN powders as raw materials. Because the reaction between Ti and Al powders is exothermic [16] , few TiN can dissolve into the Ti-Al intermetallics to form TiAl(N) solid solution[43], then few Ti2AlN particles would precipitate from TiAl(N) solid solution. Tian et al.[33] have reported that plate-like and nearly regular hexagonal Ti2AlN precipitates can form from TiAl(N) solid solution by aging after quenching from solution annealing temperatures, the plate-like and nearly regular — hexagonal Ti2AlN precipitates have a fixed orientation relationship with TiAl matrix, i.e. [112 — 0]Ti2AlN // [ 1 01]TiAl, (0006)Ti2AlN // (111)TiAl. Because both the particle morphology and crystallographic orientation relationship are same as those reported in the literature[33], the formation of hexagonal Ti2AlN particles and small plate Ti2AlN particles observed in our experiment is dominated by a dissolution-precipitation mechanism. The formation of small plate and hexagonal Ti2AlN particles occurs during the slow cooling stage of preparation process due to the reduced solubility of N atoms in TiAl matrix. As mentioned earlier, a large amount of elliptical and polygonal Ti2AlN particles have formed at the sintering temperature (1623K), so the elliptical and polygonal Ti2AlN/TiAl interfaces can act as a heterogeneous nucleation site for the precipitated Ti2AlN particles. This explains why the small plate Ti2AlN particles are more likely to attach itself to the polygonal Ti2AlN particles in our experiment. During the precipitation of the TiAl(N) solid solution, atomic matching is the dominant factor in determining the coherent interface. The atom configuration on (0001) plane of Ti2AlN and (111) plane of TiAl matrix has the best match, so the Ti2AlN(0001) plane will preferential precipitated from (111) plane of TiAl[33]. As an important member of MAX phases, Ti2AlN has a hexagonal crystalline structure with every two Ti6N octahedrals separated by an Al layer along the c-axis. In this case, the growth rate of Ti2AlN in the direction along c-axis is much lower than those parallel to the (0001) plane[44-45]. According to the above discussion, when Ti2AlN precipitated from TiAl matrix, each layer grows quickly in radial directions parallel to the (0001) plane and all the layers are stacked along [0001], then the structure ledges formed. These analysis results explains why the hexagonal Ti2AlN particles and small plate Ti2AlN particles have a ledge and terrace coherent interface with TiAl matrix in our experiment.

Conclusion 1. There are four kinds of Ti2AlN particles in the in-situ Ti2AlN/TiAl composite, which include vast

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majority of elliptical and polygonal Ti2AlN particles with the size of 4μm, few hexagonal Ti2AlN particles with the size of 4μm and small plate Ti2AlN particles with the size of 0.4μm. 2. The elliptical and polygonal Ti2AlN particles have no fixed or preferential crystallographic orientation relationship with TiAl matrix. Both the elliptical Ti2AlN/TiAl curved interfaces and polygonal Ti2AlN/TiAl flat interfaces are incoherent interphase boundaries composed by many atomic-scale microfacets. The formation mechanism of the elliptical and polygonal Ti2AlN particles is solid-state reactive diffusion, the atomic faceting in the Ti2AlN phases is the dominant factor in lowering the interface energy. 3. The hexagonal and small plate Ti2AlN particles have a fixed crystallographic orientation relationship with TiAl matrix. Both the hexagonal Ti2AlN particles and small plate Ti2AlN particles have a ledge and terrace coherent interface with TiAl matrix. The formation mechanism of hexagonal and small plate Ti2AlN particles is dissolution-precipitation, which involve a ledge growth and the lateral movement of the monoatomic layers at the interface.

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Acknowledgments

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The authors acknowledge the financial support from National Natural Science Foundation of China (grant nos. 51471058)

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GaN/α-Al2O3(0001), Physical B 324 (2002) 59-62. [37] K. Zhang, Z.J. Zhang, X.X. Lu, et al., Microstructure and composition of the grain/binder interface in WC-Ni3Al composites, Int. J. Refract. Met. Hard. Mater. 44 (2014) 88-93.

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[39] P. Li, J.M. Howe, W.T. Reynolds, Atomic structure of a {111} incoherent interface in Zr-N alloy, Acta Mater. 52 (2004) 239-248. [40] T.B. Massalski, W.A. Soffa, D.E. Laughlin, The nature and role of incoherent interphase interfaces in diffusional solid-solid phase transformations, Metall. Mater. Trans. A 37A (2006) 825-831. [41] J.F. Bartolomé, J.I. Beltrán, C.F. Gutiérrez-González, et al., Influence of ceramic-metal interface adhesion on crack growth resistance of ZrO2-Nb ceramic matrix composites, Acta Mater. 56 (2008) 3358-3366. [42] R.A. Karnesky, L. Meng, D.C. Dunand, Strengthening mechanisms in aluminum containing

ACCEPTED MANUSCRIPT coherent Al3Sc precipitates and incoherent Al2O3 dispersoids, Acta Mater. 55 (2007) 1299-1308. [43] Y. Liu, L.L. Zhang, W.W. Xiao, et al., Rapid synthesis of Ti2AlN ceramic via thermal explosion, Mater. Lett. 149 (2015) 5-7. [44] M. Magnuson, M. Mattesini, S. Li, et al., Bonding mechanism in the nitrides Ti2AlN and TiN: An experimental and theoretical investigation, Phys. Rev. B, 76 (2007) 195127. [45] G.H. Liu, K.X. Chen, H.P. Zhou, et al., Layered growth of Ti2AlC and Ti3AlC2 in combustion

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synthesis, Mater. Lett. 61 (2007) 779-784.

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Graphical abstract

ACCEPTED MANUSCRIPT Highlights •The crystallographic orientation relationships and interface atomic structures between various shaped Ti2AlN particles and TiAl matrix were studied. • The elliptical and polygonal Ti2AlN particles have no fixed crystallographic orientation relationship with TiAl matrix. • The elliptical and polygonal Ti2AlN/TiAl interfaces are incoherent interphase boundaries

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composed by many atomic-scale microfacets.

• The hexagonal and small plate Ti2AlN particles have a fixed crystallographic orientation

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relationship with TiAl matrix.

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• The hexagonal and small plate Ti2AlN particles have a ledge and terrace coherent interface with

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TiAl matrix.