Effect of N addition on microstructure refinement and high temperature mechanical properties of Ti–46Al–8Ta (at. %) intermetallic alloy

Journal Pre-proof Effect of N addition on microstructure refinement and high temperature mechanical properties of Ti–46Al–8Ta (at. %) intermetallic alloy Soroush Saeedipour, Ahmad Kermanpur, Fazlollah Sadeghi PII:

S0925-8388(19)33995-7

DOI:

https://doi.org/10.1016/j.jallcom.2019.152749

Reference:

JALCOM 152749

To appear in:

Journal of Alloys and Compounds

Received Date: 7 August 2019 Revised Date:

16 October 2019

Accepted Date: 19 October 2019

Please cite this article as: S. Saeedipour, A. Kermanpur, F. Sadeghi, Effect of N addition on microstructure refinement and high temperature mechanical properties of Ti–46Al–8Ta (at. %) intermetallic alloy, Journal of Alloys and Compounds (2019), doi: https://doi.org/10.1016/ j.jallcom.2019.152749. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Published by Elsevier B.V.

Effect of N addition on microstructure refinement and high temperature mechanical properties of Ti-46Al-8Ta (at. %) intermetallic alloy Soroush Saeedipour1, Ahmad Kermanpur1* and Fazlollah Sadeghi2 1 2

Department of Materials Engineering, Isfahan University of Technology, Isfahan 84156-83111, Iran

Pohang University of Science and Technology, Graduate Institute of Ferrous Technology, Pohang 790784, Korea

Abstract The Ti-Al-Ta intermetallic alloys are potential candidates for high temperature structural applications. This study concentrates on the microstructure refinement and high temperature mechanical properties of a Ti-46Al-8Ta alloy by N addition up to 2 (at. %). The N-bearing alloys were fabricated by vacuum are remelting, followed by hot isostatic pressing, homogenization, and solution treatments. Microstructures were characterized by XRD, OM, SEM, and TEM techniques; while mechanical properties were evaluated by hardness measurement and small punch testing. Phase transformations were determined by DTA measurements and Thermo-Calc computations. The results showed that increasing N influenced considerably morphology of the as-cast alloys, changed solidification path, and resulted in a significant microstructure refinement. The maximum small punch load at 850 °C of the fully lamellar structures enhanced at the expense of displacement at onset of fracture with increasing N content up to 1.05 at. %, beyond which the maximum load was dropped due to Ti2AlN precipitation. The hardness increased monotonically with increasing N content. The fracture surface of the low-N alloys revealed crack-tip plastic deformation, crack-bridging ligaments and plasticity around crack tips, while that of the high-N alloys showed cleavage cracks which were responsible for dramatic drop in strength.

*

Corresponding author; Tel. +98(31)33915738; Fax +98(31)33912752; E-mail: [email protected]

1

Keywords: Ti-Al-Ta intermetallic alloy; Nitrogen; Microstructure refinement; Mechanical properties; Fracture behavior; Fully lamellar structure 1. Introduction Gamma-based TiAl alloys are attractive materials for use in propulsion systems of aircraft and automotive engines due to the numerous excellent properties such as low density, good specific strength, high ignition resistance, good structural stability, and satisfactory oxidation and creep resistance at elevated temperatures [1-3]. A variety of microstructures consisting of nearly gamma, duplex, nearly lamellar and fully lamellar microstructures can be found in TiAl-based alloys [4], from which the duplex and fully lamellar microstructures were widely investigated. The duplex microstructure provides adequate ductility but poor fracture toughness and creep resistance; while the fully lamellar microstructure exhibits relatively high creep resistance and fracture toughness, but show low ductility [3, 5]. Due to the inverse relationship between ductility and fracture toughness in different microstructures, selection of final microstructure is very important to acquire optimal mechanical properties [3]. A fine-grained fully lamellar structure has been verified to possess a good combination of room-temperature ductility and high-temperature creep strength [6]. In this regard, the microstructure can be refined by powder metallurgy process [7] thermomechanical processing [8], cyclic heat treatment [9] and massive phase transformation [10]. In addition to the aforementioned methods, some studies have concentrated on the microstructure refinement by addition of interstitial elements. Schwaighofer et al. [11] showed that with increasing C content in a Ti-43.5Al4Nb-1Mo-0.1B alloy (all compositions are given in at. % unless stated otherwise), termed TNM alloy, the solidification behavior changes from the β-phase formation to a peritectic reaction. In addition, crossing the solubility limit of C led to the 2

precipitation of Ti2AlC carbides and to a refinement of lamellar colonies. A study by Tetsui et al. [12] on Ti–42Al–5Mn forged alloy has shown that increasing oxygen concentration, reduces β-phase fraction, and increases lamellar structure fraction which are required for improving high-temperature strength. Wang et al. [13] evaluated the effects of B addition on grain refinement in Ti-45Al-8Nb. Their results confirmed that two kinds of B-reduced grain refinement mechanisms through refining either β phase then α phase (β-refinement) or α phase directly (αrefinement) to refine lamellar microstructure at room temperature; however, the role of α-refinement dominated in the as-cast lamellar microstructure. Yun et al. [14, 15] studied the effect of N content on Ti-48.5Al-1.5Mo. The results indicated that with N alloying, lamellar colony size was considerably decreased in fully lamellar structure; while N causes a weak grain refinement in the duplex structure. They also demonstrated that N induces remarkable improvement of creep resistance in the duplex and fully lamellar microstructures. The Ti-46Al-8Ta alloy as a potential material for high temperature components is one of the latest generations of TiAl alloys which due to the outstanding effects of Ta in decelerating the diffusion-assisted phenomena, exhibits excellent creep and oxidation resistance [2, 10, 16]. According to the author’s knowledge, no previous work has been reported on the effect of N addition to this alloy. The aim of this investigation is to study the effect of N addition on solidification behavior, microstructure refinement, mechanical properties, and fracture behavior of fully lamellar Ti-46Al-8Ta alloy. 2. Materials and experimental procedures In this study, Ti-46Al-8Ta alloys containing different N content were prepared by vacuum arc melting from starting materials of commercially pure Ti (99.6 wt.%), Al (99.5 wt.%), Ta (99.9 wt.%) and TiN powder (99.4 wt.%) under Ar atmosphere. 3

The samples were re-melted seven times to achieve a homogeneous composition and finally, they were produced as cylinders with a diameter of 10 mm and length of 35 mm. The chemical composition of the fabricated alloys is shown in Table 1. The coding of Nxxx is used for samples in which xxx approximately represents 100×N [at%]. To minimize microsegregation and shrinkage porosity, the as-cast samples were hot-isostatic pressed (HIPed) at 1200 ℃ for 4 h at a pressure of 103 MPa and then homogenized at 1250 ℃ for 16 h, followed by furnace cooling. Solution annealing was carried out at 1350 ℃ for 2 h followed by furnace cooling to obtain a fully lamellar structure. Table 1. Chemical composition (at.%) of the studied alloys Alloy Al N000 46.23

Ta 7.84

N 0

Ti Balance

N010 46.34

7.92

0.12

Balance

N050 46.46

7.71

0.45

Balance

N100 46.19

7.68

1.05

Balance

N200 46.37

7.72

2

Balance

Differential thermal analysis (DTA) was used to determine temperature of the solid-state phase transitions. The DTA measurements were obtained by a NETZSCH STA 409 PC/PG apparatus in an Al2O3 crucible under argon atmosphere with a heating rate of 25 ℃/min from room temperature to 1400 ℃. Thermo-Calc software was also employed to calculate phase diagram of the Ti-AlTa-N system. The etching solution of 300 ml H2O, 25 ml HNO3 and 10 ml HF was utilized to reveal the microstructural features. The microstructures were investigated by optical microscopy (OM), scanning electron microscopy (SEM Philips XL30 with EDS analyzer), field-emission SEM (FEI Quanta 450) and transmission electron 4

microscopy (JEOL FE-TEM, JEM-2100F equipped with EDS analyzer and JEOL normal TEM, JEM-2100) techniques. In order to identify and validate diffraction patterns, CaRIne Crystallography v3.1 software was used. Thin foils for TEM observation were mechanically polished to about 100 µm in thickness. Then the discs were punched from the slices and jet polished using a solution of 600 ml ethanol, 360 ml butyl glycol, 60 ml perchloric acid at -10 ℃. X-ray diffraction (XRD, Philips X’Pert with Cu Kα anode) method was used for phase analysis of the solid samples. The phase percentage in optical and SEM micrographs was calculated by ImageJ software. The images were thresholded to identify and quantify each phase using a combination of ImageJ’s default thresholding algorithm and manual adjustments. Thereby, the intended phase was covered by a mask without shrinking or enlarging them. During using the ImageJ thresholding tool, the operator influence was carefully controlled, ensuring that the mask was the best possible fit for the intended phase. After applying the threshold mask, the thresholded micrograph was compared to the original micrograph to enhance accuracy in identifying the intended phase and if non-relevant indications appeared, they were removed from the measuring process. Finally, the area of each phase in the thresholded micrograph was calculated. It should be noted that this process was repeated three times to increase the measurement accuracy. Small punch test (SPT) at 850 ℃ was carried out on disc-shaped specimens that were clamped and did not move during the test. Disc slices 10 mm in diameter and 0.7 mm of thickness were cut by electro-discharge machine from the samples and then mechanically polished using the abrasive paper (P320 to P1200) to a final thickness of 0.5±0.005 mm. The force was applied by a semispherical head punch with a diameter of 2.5 mm and a constant displacement velocity of 0.5 mm/min through the test piece. This punch is driven through a receiving hole 4 mm in diameter when the specimen is deformed. All parts of the SPT device were made 5

from Nimonic-80 superalloy to provide appropriate mechanical and oxidation properties at high temperatures. During the test, the load and the punch tip displacement were continuously recorded. The punch tip was regularly inspected and if there was damage, it was replaced by a new one. In this study, the bilinear fit method was used to determine the elastic-plastic transition load (Fe) recommended by the CEN Workshop Agreement (CWA)[17]. This pre-normative document is the most recent European guidance, however, an EN standard "Metallic materials Small punch test method" is currently being formulated under the auspices of ECISS/TC101/WG11[18]. The fracture energy of SPT is determined by integration of the load punch displacement curve up to the fracture point uf: =

(Eq .1)

where uf is defined as the punch displacement at 20% load drop after maximum load. For estimating tensile properties from SP tests, the following empirical relationships (Eqs. 2-4) are used [19-22], among which Eq. 3 is more appropriate than Eq. 4 to estimate σUTS [21, 22]. =

+

= =

(Eq. 2) +

+

(Eq. 3) (Eq. 4)

Where h0 is the specimen thickness that is 0.5 mm in this study. The parameters , , , and are the calibration factors that can be calculated , empirically by comparing SP and uniaxial results. These constants rely on the small punch rig dimensions [19], sample thickness [20] and the method of extraction of Fe [19, 22]. The values for the calibration factors which have been determined by several authors on different steel grades are given in Table 2. In this regard, the optimized values by Bruchhausen et al. [19] produced satisfactory estimates for σUTS and σYS at low and high temperatures. Accordingly, in this study, due to the test rig dimensions, specimen thickness and method of extraction of Fe, their optimized factors were used to estimate the mechanical properties from 6

the SP results. Hardness of the samples was measured by Vickers method (HV) with the indenting load of 10 kg, a dwell time of 15 s and the average of 10 indentations. Table 2. The determined calibration factors in Eqs. 2-4 Calibration factor Mao et al. García et al. Altstadt et al. Bruchhausen et al. [23] [21] [20] [19] 0.360 0.476 0.48 0.382 0 0 0 28.8 0.277 0.21 0.326 0 0 -27.04 0.130 0.065 0.093 -11.86 -320 268.81 -

3. Results 3.1. Microstructures

Fig. 1 shows as-cast micrographs of the alloys with 0, 0.45 and 2 at. % N. The dendritic microstructure is mainly composed of α2/γ lamellar structure and some γ interdendritic phase. In addition to γ phase and lamellar structure, there are some β phases with bright contrast in N000 and N045 (Figs. 1d and 1e) and also some particles inside of dendrites in N200 (Figs. 1c and 1f). These particles were identified Ti2AlN by SEM-EDS analyses that exhibited rod-like or irregular morphologies. The quantitative evaluation of the as-cast micrographs by ImageJ software showed that the amount of β in N000 and N045 is 4.5 and 1 vol.%, respectively, and it is disappeared with more N alloying. On the other hand, Ti2AlN in N200 specimen was measured 1.5 vol.%.

7

Fig. 1. (a, b, c) OM and (d, e, f) SEM images of the as-cast (a, d) N000, (b, e) N050 and (c, f) N200 alloys.

A lamellar structure consists of alternating plates of γ and α2 which is formed from high-temperature α phase through a solid-state transformation in the (α+γ) phase field during cooling with a middle rate [2, 3, 24]. In order to determine the orientation relationship between two adjacent α2 and γ phases in the lamellar structure, diffraction patterns of two adjacent phases were utilized according to 8

! 0〉& are located at the Fig. 2. At a desired tilting position, 〈11!0〉$ pole and 〈112 same time. Concurrently, 0001

&

and 111

$

in both N000 and N200 alloys, are

parallel according to the diffraction patterns shown in Figs. 2c and 2f, respectively. Consequently, the following orientation relationship is derived: 0001

'

||)111*$ and < 112!0 >& || < 11!0]$

(Eq . 4)

which explains the crystallographic orientation relationship of γ lamellae with respect to α2 lamellae in one prior α grain. The obtained crystallographic orientation described in Eq. 4 is in agreement with Blackburn relationship indicating that addition of 2 at.% N does not affect orientation relationship between α2 and γ lamellar structure.

Fig. 2. TEM images of the lamellar structure in (a-c) N000 and (d-e) N200 alloys with corresponding diffraction pattern of (a, d) γ-TiAl, (b, e) α2-Ti3Al, and (c, f) the overlapped diffraction pattern of γ-TiAl and α2-Ti3Al. 9

Fig. 3 shows TEM images with corresponding selected area diffraction (SAD) pattern of a rod-like particle and the matrix in as-cast N200 alloy. The figure indicates that this particle is Ti2AlN which has formed in a γ-matrix. The nitride had no orientation relationship with TiAl-phase that results in a complete incoherent interface with the matrix.

Fig. 3. TEM images with corresponding SAD pattern of (a) the rod-like particle and (b) the matrix in as-cast N200 alloy

STEM-EDS was used to analyze the elemental distribution of constitutive phases. EDS line scan across two different particle morphologies and also the lamellar structure in N200 are shown in Figs. 4 and 5, respectively.

10

Fig. 4. STEM-EDS line scan across two different particle morphologies in N200 alloy; (a, b) cubic-like and (c, d) rod-like.

11

Fig. 5. STEM-EDS line scan across the lamellar structure in N200 alloy.

Fig. 6 indicates the calculated phase diagram of Ti-46Al-8Ta-xN and Ti-xAl-8Ta. As can be seen in Fig. 6a, N addition extends the existence range of α and σ phases, whereas reduces the β and τ phases range. Another interesting point about this figure is that Ti2N particle is more stable than Ti2AlN and Ti3AlN at low temperatures. According to these calculations, α→β+γ and β→α2+γ solid-state phase transformations occur in equilibrium conditions. Considering the fact that occurrence of α→β+γ and β→α2+γ transformations is unlikely due to the narrow temperature region of β phase, non-equilibrium solidification and also the DTA’s heating rate, some authors [25, 26], purposed to consider α→α2+γ eutectoid reaction as a replacement for these two transformations. Fig. 7a indicates the DTA graphs of the studied alloys. The measured supposed eutectoid temperature (TE) and α-transus temperature (Tα) and also the calculated value of Tα by ThermoCalc are shown in Fig. 7b. According to the calculations, N decreases Tα, but DTA results showed that the effect of N addition on TE and Tα is not uniform. 12

Fig. 6. Equilibrium phase diagram of (a) Ti-46Al-8Ta-xN and (b) Ti-xAl-8Ta, calculated by Thermo-Calc software.

13

Fig. 7. (a) DTA graphs and (b) the variation of the measured and calculated Tα and TE of the studied alloys with N content.

14

Fig. 8 shows the fully-lamellar microstructures of the studied alloys after annealing at 1350 ℃ for 2 hours followed by furnace cooling. This demonstrates that the present annealing temperature for Ti-46Al-8Ta-xN (x=0-2) locates in the single αphase field, since generating a single α-phase is essential to acquire the fullylamellar microstructure. Fig. 8f also shows EDS line scan across the lamellar structure (α2+γ) and Ti2AlN particle in N200 alloy. A sharp drop in Al and Ta concentration is seen across Ti2AlN particle, whereas Ti and N are experiencing a severe increment in concentration. Another point about the figure is that the lamellar colony size is decreased by increasing the N content. The lamellar colony size of N000, N010, N050, N100 and N200 specimens were measured 850, 550, 400, 170, and 80 µm, respectively. Fig. 9 presents the XRD profiles of the as-cast and fully lamellar alloys with different N concentrations. As can be seen, the mainly co-existing phases are α2 and γ and there is no peak of residual TiN, demonstrating complete dissolving of the TiN powders that were used as a precursor during casting. Furthermore, the observed particles were identified Ti2AlN phase. As expected, the particles appeared for the first time in more than 0.45 at. % N, which the relevant peak intensity increased with the N content.

15

Fig. 8. OM and SEM images of the fully-lamellar (a, d) N000, (b, e) N050 and (c, f) N200 alloys.

16

Fig. 9. XRD profiles of the (a) as-cast as well as (b) the fully lamellar alloys with different N concentration.

17

3.2. Mechanical properties

As γ-TiAl based alloys are regularly used as high-temperature components, SP testing was carried out at 850 ℃. Fig. 10 shows load-displacement relationship of the fully lamellar specimens with different N content at 850 ℃. It should be noted that the observed load drops before the maximum load are a consequence of crack initialization [27]. Microstructural and mechanical properties of the fully lamellar alloys that were determined from SPT at 850 ℃ and hardness test at room temperature are also indicated in Table 3. The calculated σUTS and σYS from SP results based on Eqs. 2 and 3 are shown in this table. It is worth noting that because of the transient stress state, these results are not directly comparable to those obtained from conventional uniaxial methods [28]. As can be seen in the table, with increasing N content, hardness is increased and grain size is decreased. With N alloying more than 0.45 at. %, in addition to Ti2AlN precipitation, interparticle spacing is reduced. Regarding SP results, increasing N concentration up to 1.05 at. % results in an increase in the maximum load, beyond which the load is decreased. On the other hand, N alloying decreases the displacement at the onset of fracture/failure and the fracture energy, so that the high-N alloys especially N200 exhibit poor ductility and toughness. Table 3. Mechanical and microstructural properties of the studied alloys Fe Fm ESP uf (N) (N) (J) (mm) N000 259 401 0.25 0.8

Alloy

σYS σUTS HV (MPa) (MPa) (10 kg) 395 430 297

Grain Ti2AlN λ size (µm) (Vol. %) (µm) 850 0 -

N010 315 420 0.24

0.75

481

493

325

550

0

-

N050 338 458 0.20

0.64

516

710

386

400

0

-

N100 388 498 0.17

0.45

592

811

462

170

0.5

106

N200 382 488 0.14

0.39

583

837

495

80

2

14

18

Fig. 10. Load-displacement relationship of the fully lamellar specimens with different N content at 850 ℃.

Fractography was used to evaluate the fracture mechanisms of the specimens with different N content. SEM micrographs from the fractured surfaces of SP discs after testing at 850 ℃ are depicted in Fig. 11. Fig. 12 also shows the crack path near the ruptured area of the studied alloys with fully lamellar microstructure that were tested at 850 ℃.

19

Fig. 11. SEM micrographs from the fractured surfaces of SP discs after testing at 850 ℃ for (a, b and c) N000, (d, e and f) N050 and (g, h and c) N200 alloys.

20

Fig. 12. The crack path of (a) N000, (b) N010, (c) N050, (d) N100 and (e) N200 alloys with fully lamellar microstructure that were tested at 850 ℃.

4. Discussion 4.1. Solidification path

The primary solidification phase in TiAl-based alloys strongly depends on chemical composition and casting parameters, which can be identified from the 21

dendrite morphologies. Regarding dendrite geometry during solidification, it has been demonstrated that when cubic β phase is the primary phase, it develops in three orthogonal <100> directions and presents a 4-fold symmetry. However, when hexagonal α phase is the primary phase, it develops along <101!0> directions in the basal plane, and one more direction orthogonal to these along the c-axis [0001] and presents a 5 or 6-fold symmetry [29-31]. As shown in Fig. 1, the angle between the dendritic arms for N000 and N010 alloys was approximately 90°, while it was approximately 60° for N100 and N200 alloys. Therefore, it can be perceived that β has formed as the primary solidification phase in the low-N alloys (Fig. 1a) and α as the primary solidification phase in the high-N alloys (Fig. 1c). For the intermediate N content, two symmetry characteristics were observed (Fig.1 b). The angle between the dendritic arms in some regions was about 90° and in other regions, it was about 70°. Indeed, the first sign of formation of α as the primary solidification phase was appeared in 0.45 at. % N. The interesting point about solidification path of the studied alloys is that according to the calculated phase diagram of Ti-46Al-8Ta-xN (Fig. 6a), the primary solidification phase for Ti-46Al-8Ta-xN (x=0 to 2) is β. Whereas, the geometry characteristics of N100 and N200 dendrites show that solidification has occurred through α primary phase (Fig. 1). The most effective reasons for changing the primary phase in High-N alloys will be discussed in the following section. From the thermodynamic point of view, N content exceeds the solubility limit in N100 and N200 samples and as a result, the formation of Ti2AlN is occurred. Precipitation of Ti2AlN causes Ti depletion of the adjacent area and accordingly, Al/Ti ratio is increased which is in favor of the formation of α as the primary solidification phase. In addition, N is an α-stabilizer element that extends α field. Therefore, the possibility of α phase appearance as primary solidification phase in 22

the high-N alloys is more than those with the lower N content. As can be seen in Fig. 6a, increasing N content makes the β stability region narrower. In addition, whereas the calculated phase diagram (Fig. 6a) corresponds to equilibrium conditions, due to the use of a water-cooled copper crucible, the liquid experienced a high cooling rate resulting in non-equilibrium solidification. Consequently, for the high-N alloys under non-equilibrium conditions, the possibility of β phase appearance as the primary solidification phase is severely reduced. 4.2. Segregation pattern

The unexpected point about the chemical profile across the nitrides (Fig. 4) is that in a narrow band with 40-50 nm width (marked as *) around both types of Ti2AlN, Ti concentration is increased abruptly and then decreased. The composition of this band is very similar to α2 (Ti3Al). Since Ti2AlN is the primary solidification phase in the high-N alloys, one possible hypothesis is that the high N content of α phase around the primary Ti2AlN at the temperatures near to Tα, stabilizes α phase, such that this phase remains untransformed during cooling down to room temperature. In addition, γ precipitation from α phase is a diffusioncontrolled transformation and it is suggested that N slows down diffusion rate of Ti and Al atoms [14]. On the other hand, diffusion is occurred in only one side of the Ti-rich band, while the other side is in contact with Ti2AlN particles. Distribution of the alloying elements across the lamellar structure is indicated by the line scan profile of STEM-EDS analysis (Fig. 5). As can be seen, there is a difference between Ti and Al content across the lamellar structure, so the laths with dark and grey contrast identified γ and α2, respectively. In addition, there is no element segregation along lamellar interfaces. Considering the fact that the system energy is reduced especially when an element segregates to an incoherent or semicoherent interface, it was already expected that alloying elements were enriched in 23

these regions. This is maybe due to the limitation of the characterization used to analyze such interfaces with about 15 nm in thickness. Another interesting point is that N and Ta localized in the α2 lath. This finding is in agreement with the study by Kainuma et al. [7] that showed Ta always partitions to α or α2 more than to γ. On the other hand, the higher N solubility in α2 phase (D019-structure) contrast to γ phase (L10-structure) is originated from their octahedral cavities. These cavities are Al4Ti2 and Al2Ti4 types with the same number of each type in L10-structure which are not appropriate sites for N atoms. On the other hand, octahedral cavities of the D019-structure are Al2Ti4 and Ti6 types. Ti6 cavity is a suitable site for N atoms which results in an increase of N solubility in α2 phase in contrast to γ phase [32]. In this study, the ratio of N and Ta content in α2 lath to that in the γ lath is about 1.5 and 1.3, respectively. It seems that the partitioning degree of interstitial atoms is more than that for substitutional atoms, which is in agreement with the previous study [33]. 4.3. Critical temperatures

N is a powerful α-stabilizer element that widens α phase field and shifts α phase stability to lower temperatures. Therefore, as can be seen in the calculated phase diagram (Fig. 6a), Tα should be decreased by N alloying. However, DTA measurements showed that N addition beyond 0.45 at. % increases Tα (Fig. 7). This contradiction is originated from the fact that the N solubility limit of Ti-46Al-8Ta is approximately around 0.45 at. % and a further increase in N content would result in precipitation of the coarse Ti2AlN particles. Although N acts as an α-stabilizer element and decreases Tα, however, Ti concentration of matrix is depleted as a result of nitride precipitation during solidification beyond 0.45 at. % N and consequently, Al/Ti ratio is increased. As can be clearly seen in Fig. 6b, increasing

24

the Al/Ti concentration, increases α-tansus temperature. This would explain increasing of Tα when N addition is more than 0.45 at. %. Concerning eutectoid temperature (TE), N addition up to 0.45 at. % increases TE, while further N alloying does not have a significant effect. It seems that the dissolved N atoms like C stabilize α2 and increase TE [34, 35]; On the other hand, when N is in concentrations exceeding saturation limit, it is expected to destabilize of α2 due to the increase in Al/Ti ratio, which limits the large increment in TE [36]. 4.4. High temperature mechanical properties

It is interesting to note that N100 alloy has the highest max load in the SP testing while having lower Ti2AlN volume fraction, grain size, interparticle spacing and even hardness in contrast to N200 alloy (Table 3). The most effective mechanisms in reducing the max load of the high-N alloys will be discussed in the following section. The increment of the strength up to 0.45 at. % N is originated from solid solution strengthening and grain-refining effect of solute N atoms. With further N alloying, nitrides precipitation can occur affecting strength directly or indirectly. The nitrides induce direct strengthening by interaction with dislocations and increasing dislocation density around Ti2AlN/matrix interface during deformation. In this regard, it has been proposed that the strength is increased linearly with λ-1/2 where λ is interparticle spacing. It should be noted that in this study the interparticle spacing values were too big to have a remarkable effect on the strength via Orowan mechanism. On the other hand, the indirect strengthening effect originates from grain refinement due to grain boundary pinning of Ti2AlN precipitates. In this case, yield strength is increased linearly with D-1/2 where D is grain size [2, 37, 38]. The average max load at 850 ℃, the hardness at room temperature, the value of λ-1/2 (direct strengthening) and D-1/2 (indirect strengthening) as a function of N content 25

are plotted in Fig. 13. Due to the increasing of λ-1/2 and D-1/2 with N content, it was expected that the max load similar to the hardness value would have a direct relation with N content. However, surprisingly the specimen with 2% N fractured at a load even lower than those for N100 alloy. It seems that the deterioration of the max load in N200 alloy is related to the higher amount of accumulated nitrides which act as crack initiation sites.

Fig. 13. The max load at 850 ℃, hardness at room temperature, the value of λ-1/2 and D-1/2 as a function of N content.

4.5. Fracture behavior

As expected, macroscopic fracture surfaces of the investigated alloys (Figs. 11a, 11d and 11g) indicated brittle fracture mechanism. However, N200 alloy even showed more brittle appearance than those alloys with the lower N content. As marked in Fig. 11g, there are signs of circumferential cracks in addition to the radial cracks in the fracture surface of this alloy. Viewing the microscopic fracture surfaces of the tested alloys at 850 ℃ reveals mixed-mode failures, involving both 26

of interlamellar and translamellar fracture modes (Figs.11b, 11e and 11h). Regarding microstructural features, the N-free alloy shows a larger interlamellar spacing. With increasing N concentration, a rougher fracture surface can be generally observed demonstrating the effect of N content and grain refinement on the fracture morphology. Another point is that the fracture surface of N000 alloy (Fig.11b) presents evidence of an interesting crack path through the lamellar structure. As seen in the figure, the crack is propagated translamellarly along few lamellae, then deflected interlamellarly for a limited extent, and this is repeated for several times. It seems that this behavior is originated from the influence of lamellar orientation on the crack path that will be discussed later. At higher magnifications (Figs. 11c, 11f and 11i), some features such as delamination along α2/γ interfaces, translamellar and interlamellar fracture are clearly revealed. About interlamellar fracture, it should be noted that the flat planes and irregular polygonal represent cleavage-like fracture which could be originated from parallel interlamellar microcracks [39]. In addition, in the high-N alloy (Fig.11i), backscattered electron (BSE) imaging has been used to evaluate the existing phases on the fracture surface. As it is marked on the micrograph, the phase with darker contrast is Ti2AlN that was proved by EDS analysis. The fact that in the N200 alloy there are some areas with a high volume fraction of Ti2AlN nitrides, suggests that stress concentration is possible to be relaxed by the initiation of cleavage cracks since stress relaxation by crack nucleation will be facilitated more than plastic flow [40]. Consequently, it leads to dramatically drop in the strength and fracture toughness as earlier shown by SP results (Fig. 10 and Table 3). Concerning the crack profile (Fig. 12), crack propagation occurred by a mixed mode of translamellar, interlamellar and even intergranular fracture through the lamellar colonies in all alloys. Clearly, in contrast with that observed in N200 alloy 27

(Fig. 12e), crack-tip shielding effect has occurred in the lower N content alloys (Figs. 12a-12d). Fig.14 shows a blunted crack at higher magnification in N000 specimen. As can be seen, there are some clear evidence of the lamellae deformation around the blunted tip. In these alloys, crack tips will be blunted when interacting with lamellae. As the load is increased, the crack re-nucleated in front of the blunted tips, consequently intact ligament bridges two separated sections of the crack [41]. Mercer et al. [41] proposed that the interrupted crack growth and consequently the formation of intact crack-bridging ligaments and then crack renucleation in front of the blunted cracks is a consequence of interactions of the crack tips with the α2 lamellar boundaries.

Fig. 14. Crack tip shielding in N000 alloy

Lamellar orientation has a great effect on fracture behavior. It has been demonstrated that increasing the angle between the lamellar plane and crack plane 28

results in an improvement of fracture toughness [2, 39, 42, 43]. Wang et al. [39] investigated the effect of lamellar orientation on the fracture behaviors of Ti-48Al2Cr-2Nb (at.%) poly-synthetically twinned crystals. They showed that the fracture toughness increases with the increased oriented angle, and the fracture behavior changes from interlamellar to translamellar. Fig. 15 shows effect of inclination angle between the crack plane and lamellae orientation on fracture behavior of the present N010 alloy. The stress arrows shown in Figs. 15d, 15e and 15f were used to simplify of the fracture mechanism. It is worth noting that the loading configuration in the SP test is predominantly biaxial stress state and differs from uniaxial loading in a standard tensile or creep rupture test [17]. As can be seen in Figs. 15a and 15d, there are a few ligaments that are broken during the load increasing. When the inclination angle between the crack plane and lamellae orientation is increased, fracture mode is still interlamellar but it is toughened by plasticity and creating microcrack and ligament (Figs. 15b and 15e). The ligaments would increase fracture toughness since more energy is needed to tear ligaments between the main crack and microcracks. In addition, a deformed area can be seen in Fig. 15b, indicating plasticity is enhanced by increasing the inclination angle between the crack plane and the lamellae orientation. When the crack plane is almost perpendicular to lamellae orientation (Figs. 15c and 15f), translamellar fracture mode is predominant. In this case, plasticity and fracture toughness would be maximum. This is due to the fact that as the main crack advances through the material and encounters with a lamellae interface, there are some differences in the structure and the orientation between the neighboring α2 and γ lamellae. On the other hand, plasticity around the crack can induce blunting its tip; accordingly, the crack is forced to stop. As the load is increased, the crack re-nucleates along a new cleavage plane and propagates until it stops by next barrier at the α2 and γ interface. This trend is repeated several times and finally, the crack travels into the 29

structure through a zig-zag path. Crack blunting, renewing and deflection processes take a lot of energy to proceed which is in favor of fracture toughness.

Fig. 15. The crack interaction with lamellar structure when (a, d) the lamellae are parallel to crack plane, (b, e) inclination angle between the crack plane and the lamellae is increased and (c, f) the lamellae are perpendicular to crack plane in N010 alloy.

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5. Conclusions

This study concentrated on microstructure refinement and high temperature mechanical properties of N-bearing Ti-46Al-8Ta (at. %) alloys. The main conclusions are summarized below. 1. γ and α2 lamellae had Blackburn orientation relationship with each other in the lamellar structure and addition of up to 2 at. % N did not have any influence on this orientation relationship. 2. N alloying in excess of the solubility limit (nearly 0.45 at. %) led to the Ti2AlN precipitation with no sign of the formation of Ti3AlN particles. 3. N alloying resulted in dendrite refinement and decreasing the retained metastable B2 phase in the as-cast state. The grain size in the fully-lamellar specimens was also decreased with increasing N content. 4. Increasing N content influenced considerably morphology of the as-cast samples and changed the β primary phase to the α phase during solidification. 5. Increasing N content up to 1.05 at. % resulted in an increase in the maximum small punch load at the expense of displacement at the onset of fracture and also the fracture energy in the fully lamellar structure. However, further N alloying despite grain refinement and Ti2AlN precipitation led to a drop in the maximum load. 6. The hardness increased monotonically with increasing N content. 7. The fractography evaluation showed that cracks propagate through the fully lamellar structures with different N content by a mixed mode of translamellar, interlamellar and even intergranular manner. 31

8. While crack-tip plastic deformation, crack-bridging ligaments and plasticity around crack tips occurred in the fracture surface of the low-N alloys, a high amount of accumulated nitrides act as crack initiation sites and deteriorates mechanical properties of the high-N alloy. 9. When the inclination angle between the crack plane and lamellae orientation is increased, crack growth will be more difficult.

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Figure captions: Fig. 1. (a, b, c) OM and (d, e, f) SEM images of the as-cast (a, d) N000, (b, e) N050 and (c, f) N200 alloys. Fig. 2. TEM images of the lamellar structure in (a-c) N000 and (d-e) N200 alloys with corresponding diffraction pattern of (a, d) γ-TiAl, (b, e) α2-Ti3Al, and (c, f) the overlapped diffraction pattern of γ-TiAl and α2-Ti3Al. Fig. 3. TEM images with corresponding SAD pattern of (a) the rod-like particle and (b) the matrix in as-cast N200 alloy. Fig. 4. STEM-EDS line scan across two different particle morphologies in N200 alloy; (a, b) cubic-like and (c, d) rod-like. Fig. 5. STEM-EDS line scan across the lamellar structure in N200 alloy. Fig. 6. Equilibrium phase diagram of (a) Ti-46Al-8Ta-xN and (b) Ti-xAl-8Ta, calculated by Thermo-Calc software. Fig. 7. (a) DTA graphs and (b) the variation of the measured and calculated Tα and TE of the studied alloys with N content. Fig. 8. OM and SEM images of the fully-lamellar (a, d) N000, (b, e) N050 and (c, f) N200 alloys. Fig. 9. XRD profiles of the (a) as-cast as well as (b) the fully lamellar alloys with different N concentration. Fig. 10. Load-displacement relationship of the fully lamellar specimens with different N content at 850 ℃. Fig. 11. SEM micrographs from the fractured surfaces of SP discs after testing at 850 ℃ for (a, b and c) N000, (d, e and f) N050 and (g, h and c) N200 alloys. Fig. 12. The crack path of (a) N000, (b) N010, (c) N050, (d) N100 and (e) N200 alloys with fully lamellar microstructure that were tested at 850 ℃. Fig. 13. The max load at 850 ℃, hardness at room temperature, the value of λ-1/2 and D-1/2 as a function of N content. Fig. 14. Crack tip shielding in N000 alloy. Fig. 15. The crack interaction with lamellar structure when (a, d) the lamellae are parallel to crack plane, (b, e) inclination angle between the crack plane and the lamellae is increased and (c, f) the lamellae are perpendicular to crack plane in N010 alloy.

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Table captions: Table 1. Chemical composition (at.%) of the studied alloys Table 2. The determined calibration factors in Eqs. 2-4 Table 3. Mechanical and microstructural properties of the studied alloys

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Highlights • Effect of N on solidification behavior and microstructure refinement is evaluated. • Effect of N on HT mechanical properties and fracture behavior is investigated. • Increasing N resulted in a significant microstructure refinement. • Ti-46Al-8Ta alloy containing about 1 at.% N showed the maximum strength. • Crack propagation occurred by a mixed mode of translamellar and interlamellar. • Inclination angle between crack plane and lamellae orientation affected crack path.

Tuesday, 22 October 2019

Professor Jennifer Ann Aitken, The Editor, Journal of Alloys and Compounds, Elsevier RE: Effect of N addition on microstructure refinement and high temperature mechanical properties of Ti-46Al-8Ta (at. %) intermetallic alloy (Manuscript ID: JALCOM-D-19-10509) Dear Professor Aitken, This is to confirm that the work described in the above mentioned manuscript has not been published previously, that it is not under consideration for publication elsewhere, that its publication is approved by all authors and that, if accepted, it will not be published elsewhere in the same form, in English or in any other language, without the written consent of the publisher. Please inform us if any further information are needed. Look forward to hearing from you. Yours Sincerely,

Ahmad Kermanpur

Soroush Saeedipour

Fazlollah Sadeghi

Department of Materials Engineering Isfahan University of Technology Isfahan 84156-83111, Iran Tel: +98 (0) 31 3391 5738 Fax: +98 (0) 31 3391 2752 Mobile: +98 (0) 913 325 6705 E-mail: [email protected]