Effects of HIP on microstructural heterogeneity, defect distribution and mechanical properties of additively manufactured EBM Ti-48Al-2Cr-2Nb

Accepted Manuscript Effects of HIP on microstructural heterogeneity, defect distribution and mechanical properties of additively manufactured EBM Ti-48Al-2Cr-2Nb Mohsen Seifi, Ayman A. Salem, Dan P. Satko, Ulf Ackelid, S. Lee Semiatin, John J. Lewandowski PII:

S0925-8388(17)33212-7

DOI:

10.1016/j.jallcom.2017.09.163

Reference:

JALCOM 43219

To appear in:

Journal of Alloys and Compounds

Received Date: 23 November 2016 Revised Date:

14 September 2017

Accepted Date: 16 September 2017

Please cite this article as: M. Seifi, A.A. Salem, D.P. Satko, U. Ackelid, S.L. Semiatin, J.J. Lewandowski, Effects of HIP on microstructural heterogeneity, defect distribution and mechanical properties of additively manufactured EBM Ti-48Al-2Cr-2Nb, Journal of Alloys and Compounds (2017), doi: 10.1016/ j.jallcom.2017.09.163. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. 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.

ACCEPTED MANUSCRIPT

Effects of HIP on Microstructural Heterogeneity, Defect Distribution and Mechanical Properties of Additively Manufactured EBM Ti-48Al-2Cr-2Nb Mohsen Seifi1, *, Ayman A. Salem2, Dan P. Satko2,

RI PT

Ulf Ackelid3, S. Lee Semiatin4, and John J. Lewandowski1

1

Department of Materials Science and Engineering Case Western Reserve University, Cleveland, OH 44106, USA

2

Materials Resources LLC, Dayton, OH 45402, USA

3

Arcam AB, Molndal; SE-431 37, Sweden

4

[email protected]

M AN U

*

SC

Air Force Research Laboratory, Materials and Manufacturing Directorate, AFRL/RXCM, Wright-Patterson Air Force Base, OH 45433, USA

Abstract

The present work was conducted to establish an initial understanding of the processingmicrostructure-property relations for γ-TiAl manufactured by electron beam powder bed fusion. The investigation included microstructure characterization at different length scales of near-γ titanium aluminide Ti-48Al-2Cr-2Nb (in atom pct) in addition to evaluating the compressive

TE D

strength, microstructures, fracture toughness, and fatigue crack growth behavior. Micro-CT revealed significant variations in the spatial distribution of internal defects along the build of the as-deposited sample, while HIP reduced these defects. However, inhomogeneous microstructures were exhibited in both the as-deposited and HIP materials while HIP processing reduced both the

EP

yield strength and amount of scatter in mechanical properties. Despite these observations, fracture toughness and fatigue crack growth results were in the range of those reported for

AC C

conventionally cast material. The use of cloud computing provided an efficient means for data management, microstructure analytics, and collaboration among various contributors to the work.

Keywords: Additive Manufacturing, Gamma Titanium aluminide, Ti-48Al-2Cr-2Nb, Electron beam melting,

microstructure evolution, Defect, Tomography, fracture toughness, fatigue

1|Page

ACCEPTED MANUSCRIPT

1. Introduction Titanium aluminides (TiAl) possess low density, high specific strength/modulus at

RI PT

elevated temperatures, as well as excellent oxidation and creep resistance, particularly in comparison to nickel-base superalloys [1]–[14]. While these properties are highly desirable for next generation aerospace and automotive applications, only recently (i.e. in 2010) have cast Ti48Al-2Cr-2Nb blades been certified and implemented in GE commercial turbofan engines [12],

SC

[15]–[17]. The casting of γ-TiAl alloys has some processing challenges in addition to machining difficulties that arise from the low ductility inherent with many intermetallic phases [1], [18]–

M AN U

[21]. The limited number of slip systems in γ-TiAl can also produce relatively low fracture toughness [22]–[24].

Attempts to explore other processing routes include a number of different freeform fabrication techniques to produce TiAl structures to near net shape [25]. One approach provides

TE D

the possibility of producing functionally graded alloys by using gas tungsten arc welding (GTAW) as the heat source [26], [27]. Other approaches include both laser-based [22], [28]–[32] and electron-beam-melting (EBM) techniques [28], [33]–[52] to melt pre-alloyed powders to

EP

produce TiAl components. The latter also provides the opportunity to produce parts with lower

AC C

residual stress due to the pre-heated powder bed and high-temperature processing, as well as low oxidation because of the high vacuum environment. However, the high vacuum environment can produce preferential vaporization of Al and Cr [45]. Despite these complications, EBM has a deposition rate that is higher in comparison to laser-based techniques and is being explored to produce complex, near net shape γ-TiAl parts with short lead times for various applications [12], [28], [33]–[50]. Recently, it was reported that GE aviation is considering the implementation of EBM TiAl to replace cast TiAl [53]. 2|Page

ACCEPTED MANUSCRIPT

Due to the nature of the process and relatively rapid directional solidification of a localized melt pool in EBM powder bed fusion, the microstructure and resulting mechanical properties may differ substantially from those obtained with conventional processing techniques.

RI PT

In the case of TiAl, alloys considered for structural applications typically contain two main phases, γ-TiAl as the matrix and α2-Ti3Al with various morphologies and spatial distributions depending on the type of processing/heat treatment. The present work was conducted to

SC

characterize the microstructure, compressive strength, defect population, fracture toughness, and fatigue crack growth behavior of electron beam powder bed AM γ-TiAl in the as-deposited and

M AN U

HIP conditions. Very limited work has investigated the fracture and fatigue behavior [49] of EBM γ-TiAl alloys. It was suggested [47], [49], [51] that one advantage of EBM processing of γTiAl relates to the potential of avoiding/minimizing defects typically present in cast or powder metallurgy (PM) techniques.

TE D

2. Materials and Methods

An Arcam 3kW EBM machine, model A2X, was used to produce 15 mm diameter × 75

EP

mm length multi-layer circular rods, Figure 1(a), using Arcam pedigree gas-atomized (from the pre-alloyed melt) Ti-48Al-2Cr-2Nb spherical powders with an average particle size of 45-150

AC C

µm. The elemental powder chemistry is summarized in Table 1 with a comparison to the alloy specification for Ti4822. Standard Arcam EBM zigzag raster strategy operated at about 1050 °C process temperature with layer thickness of 90 µm was used to deposit cylindrical bars at 10-5 mbar chamber pressure. Each powder layer was pre-heated at 1050 °C by employing a scanning beam current of 10 mA and scanning velocity of 5 m/s. Some material was HIPped at 1200 °C (+/-10 °C) (i.e. in the two-phase, α + γ field) and 103.4 MPa for 4 hr in an effort to close any porosity and attempt to homogenize the microstructure. This initial selection of 1200 °C was 3|Page

ACCEPTED MANUSCRIPT

inspired by a number of previous works [15], [38], [45], [50], [51] on both conventionally cast and AM materials, while HIP at 1280 °C was shown to remove porosity in cast ingots of Ti46Al-8Ta [54]. However, the present work suggests that heat treatment/HIP in the single-phase

RI PT

alpha field will likely be required to eliminate micro-segregation, as discussed later. Both asdeposited and HIPped specimens were subsequently machined to 10 mm × 10 mm × 75 mm bend bars, Figure 1(b).

SC

Detailed microstructure characterization was conducted midway from the start to the end

M AN U

of the build (side view) in both the as-deposited and HIPped conditions, Figure 1(a) and Figure 2, at Materials Resources LLC (MRL) using a TESCAN scanning electron microscope (SEM) equipped with electron backscatter diffraction (EBSD) detector. The samples were prepared for EBSD using standard metallographic methods. Backscattered electron (BSE) imaging and large area EBSD scans were conducted to investigate the spatial distribution of microstructure

TE D

features. Large area EBSD data were collected using an in-house beam/stage control algorithm that allowed for automatic EBSD scans of 110 scans for the as-deposited conditions and 90 scans for the as-HIP conditions (each 240 µm x 240 µm) with a step size with 0.8 micrometers. For

EP

limited scans of the HIP material, both γ (TiAl) and α2 phases were collected for phase

AC C

segmentation. Data analytics, data management, and sharing among team members was conducted on MiCloud [55] which was also used for statistical analysis of defects that was measured by µCT scans. In addition, EDS analyses of 2.5 mm × 2.5 mm areas were conducted on polished samples taken at the start of a build, at the middle (37.5 mm from the start of the build) of the same build and at the ¼ position (18.5 mm from the start of the build). Compression tests were conducted on an MTS© servo-hydraulic machine at a strain rate of 2 × 10-4 s-1 on sub-sized cylindrical specimens shown in Figure 1(d) to estimate the 4|Page

ACCEPTED MANUSCRIPT

compressive strength and flow behavior. Fracture toughness testing was performed initially on the as-deposited cylindrical specimens shown in Figure 1(a) and Figure 2 using the stress intensity factor calculation for three- point bending derived from [56]. Subsequently, single edge

RI PT

notch (SEN) specimens taken from the machined 10 mm x 10 mm x 75 mm samples shown in Figure 1(b), (c) and Figure 2 were tested in three-point bending on a Model 810 MTS servohydraulic machine in general accordance with ASTM E399 [57]. An attempt was also made to

SC

obtain fatigue crack growth data according to ASTM E647 [58]. In both cases, the SEN specimens were first notched using a slow speed diamond wire saw to introduce a notch with

M AN U

root radius of about 100 µm. Fatigue pre-cracking was not successful due to the difficulty in initiating a stable fatigue crack. Compression fatigue pre-cracking in four point bending based on the recommendations of [59]–[63] also was not successful. For this reason, fatigue crack growth testing was performed using incremental load increases to promote crack initiation at low ∆K

TE D

using constant cyclic loading followed by dcPD (direct current potential drop) monitoring of crack growth under rising ∆K conditions to failure explained elsewhere [64]. Values for Paris slope and Kmax at failure were recorded along with the ∆K at which fatigue cracking was

EP

initiated. Since this is not the standard method for determining ∆Kth, it will be designated as

AC C

∆K*th in the present work to indicate this non-standard estimation of the fatigue threshold. X-ray micro-computed tomography (µ-CT) was carried out at YXLON (a division of

Comet technologies) using the YXLON FF35 system at 195 kV and 120 µA with a resolution of about 20-30 µm. Scans were performed on a YXLON FF35 system using a helical scanning technique to scan the whole sample instead of using a multiple-section scan. Three-dimensional defect volumes were determined by rotating the sample 360°, enabling the detector to capture numerous two-dimensional projected images from which a three-dimensional data set of volume

5|Page

ACCEPTED MANUSCRIPT

elements (voxels) was reconstructed based on appropriate algorithms and high computing power computers using VG Studio MAX 2.2 software [65].

RI PT

3. Results Microstructure examination of the as-deposited material, Figure 3, revealed a duplex structure consisting two regions: a single-phase region and a two-phase region. The single-phase

SC

region is characterized with coarse equiaxed γ-grains. The two-phase regions are characterized by fine γ grains and α2 phases. High resolution BSE images, Figure 3 (c), shows the presence of

M AN U

limited areas (less than 20 µm wide) with fine lamellar features and fine equiaxed γ grains in the as-built material. The as-deposited microstructure evolves from the double cascading peritectic. The α2 + γ have evolved from the initial dendritic solidification product. The (interdendritic) γ grains form from the second peritectic reaction (α + L
γ). Thus, no α /α2 would be expected

TE D

within the single-phase γ region. The fact that such microsegregation has been observed previously [66], [67] in TiAl synthesized by both PM techniques and ingot-metallurgy techniques suggests that the microsegregation is not highly dependent on cooling rate. However,

EP

the length scale of the microsegregation is indeed a function of the cooling rate. Figure 4 confirms the microsegregation after HIP process by segmenting the EBSD data collected from

AC C

the two-phase region (with fine γ grains and α2 phases). In particular, the build direction inverse pole figure maps (Figure 4a and b) revealed a 19% and 81% volume fraction of α2 and γ, respectively. Texture analysis by pole figures for each phase (Figure 4c and 4d) revealed intermediate texture intensity for each phase (4 and 3.4 times random for α2 and γ, respectively) with almost all of the [001]γ poles were aligned perpendicular to the build direction (RD in Figure 4c-4d) and the [0001]α2 poles were split between ~20°-35° tilt from the build direction and 90° to the build direction. 6|Page

ACCEPTED MANUSCRIPT

The as-deposited material also contained a limited number of microcracks and voids as shown in Figure 3. Although these will be detrimental to the mechanical properties in certain test orientations, it has been reported that HIP can be used to remove such defects [38], [39].

RI PT

However, high-resolution imaging of the HIPped material, Figure 5, shows the presence of a micro-crack within a single γ grain along with unusual crack deflection across the matrix/twin boundaries within that single γ grain. Also, secondary twins within primary twins were noted,

SC

Figure 5(a). Both low and high magnification BSE revealed the presence of microstructure heterogeneity after HIP, Figure 5(b) and (c), with layered bands of small γ grains within areas of

M AN U

large γ grains. It is worth noting that the lamellar regions disappeared after HIP treatment (Figure 5). Previous studies [68] have shown that the typical grain size for as-cast γ-TiAl is 1000 µm ± 140 µm. In contrast, the EBM as-deposited material, Figure 3 exhibits an average grain size of only 10-40 µm. This observation is similar to the other orientation (top view) reported

TE D

previously [69]. This significant microstructure refinement is responsible for the elevated hardness/strength reported by others in similarly processed materials [38], [39]. In order to obtain a broader view of the microstructure, large area EBSD was used to

EP

record the microstructure using 110 scans for the as-deposited conditions and 90 scans for the as-

AC C

HIP conditions (each 240 µm x 240 µm). These scans were analyzed using Ti-ZoneTM module on MiCloud.AM [70] forming the montage shown in Figure 6 for as-deposited material that reveals the spatial heterogeneity of the microstructure. The use of cloud computing [55] allowed efficient data sharing among team members and efficient data analysis of the large number of EBSD scans by uploading the dataset once and then team members were able to analyze and visualize the results using thin or thick client platforms (e.g. tablet, laptops, etc) without the need to download or install data or software. The online data analysis by TiZone included the plotting

7|Page

ACCEPTED MANUSCRIPT

of inverse pole figure maps and the calculations of orientation distribution function using generalized spherical harmonics (GSH). These were conducted on individual scans and then the final results were assembled together using TiZone GSH algorithm [71]. Similarly, microCT and

RI PT

fracture toughness results were uploaded to MiCloud and the statistical analysis by boxplots was carried out using the Analytics web-app. This enabled more efficient discussion of results be inspecting the same output live on the internet without being in the same state.

SC

In particular, for as-deposited material, the light gray ribbon-like looking areas in Figure

M AN U

6(a) were confirmed at higher magnification to be regions of coarse γ grains, Figure 6 (b) and (c). Segmenting the dataset based on the median grain size produced two distinctly different subsets consisting of coarse and fine γ grain as-deposited material, Figure 7. Statistical representation of the grain size distribution using the box and whisker plot, Figure 7 (d), show that the median grain size of the coarse grains (21.9 µm) was almost 6 times that of the fine grain

TE D

size (3.9 µm) in the as-deposited material. The presence of annealing twins was only detected within coarse γ-grains. This microstructure heterogeneity occurred within a uniform shaped (i.e. cylindrical) part that was built using spatially uniform processing parameters (i.e. along the build

EP

direction). This reflects the material’s sensitivity to the solidification process and any location-

AC C

specific temperature profiles along with contributions from chemical segregation [66], [72]. This microstructure heterogeneity is expected to have a direct impact on the location-specific mechanical properties. Nano-indentation conducted in the fine and coarse regimes using a Nano Indenter model G200 made by Keysight (formerly Agilent) operated under depth control of 1000 µm revealed 4.34 GPa and 3.61 GPa respectively. EDS analyses, Figure 8, shows a variability in the chemical element concentration as measured by EDS in the single phase and two- phase bands of HIPped sample. The average Al 8|Page

ACCEPTED MANUSCRIPT

concentration varies between 44.18 – 46.15 at pct. The two-phase bands showed lower Al concentration than the single-phase bands. The small difference in the average Al concentration between the single phase and the two-phase bands is due to the large volume fraction of the γ

RI PT

grains in the two-phase regions compared to the α2 phase. In particular, the γ in the single phase/large γ grain region must have the same composition as the γ in the duplex (α2 + γ) fine grain region, assuming we are close to equilibrium. Hence the small differences captured in

SC

Figure 8 are the result of probing single phase regions which are almost entirely γ (spectra 1, 3, 5) whose average Al content would be on the high side vs probing two-phase regions which are a

M AN U

mixture of high Al content γ (at high volume fraction) and lower Al content α2 (at lower volume fraction) (spectra 2, 4). It is worth noting that the data in Figure 8 suggest that the HIP treated material had Al concentration of 45-46 at% not 48 at%. The source(s) of the drop in the Al concentration with respect to original powder will be investigated further in future work.

TE D

The spatial variation of the microstructure along with the use of HIPping as a mitigation process to close pores and reduce defects requires the use of large area EBSD in order to characterize the as-deposited microstructure in addition to establishing the impact of any post

EP

processing. To minimize the impact of location of the inspection area, the same spatial location

AC C

was selected in both samples by EBSD. That is, the same distance from the outside sample surface in x, y, and z directions. The image quality map from 110 EBSD tiles (each 240 µm x 240 µm) obtained before and after HIPping, Figure 9, revealed the presence of bands of coarse and fine γ grains.

Quantitative statistical analysis of the HIPped material, similar to that

conducted for the as-deposited material, Figure 7, also revealed the persistence of banded coarse and fine γ grains, Figure 10. The box and whiskers plot in Figure 10 revealed that the coarsegrained regions possessed grain sizes almost 6 times that of the fine-grained regions (e.g. 20.4

9|Page

ACCEPTED MANUSCRIPT

µm vs. 4.4 µm). Furthermore, quantitative comparison of the grain size statistics in the asdeposited and HIP conditions, Figure 11, confirmed that HIP produced some changes in the equiaxed grain size and microstructure heterogeneity in both single phase and two-phase regions.

RI PT

However, Figure 3 and Figure 5 showed that HIP has eliminated the lamellar colonies (perhaps via a spheroidization/termination migration process) as well as led to a general coarsening of both the α2 and fine γ grains.

SC

Spheroidization and coarsening should reduce the Hall-Petch contribution to strength,

M AN U

although it was also shown that there was only a limited change in grain size in the single phase γ grain bands (Figure 11). Since the metallographic analyses have shown that these single phase γ grain bands occupy approximately 50% of the overall microstructure, the lack of grain size change within the single phase γ grain bands should minimize the impact of the spheroidization/coarsening in other regions on the global strength, assuming the strength of such

TE D

structures follows a rule of mixtures.

These results also confirm the need for large area characterization (100s of EBSD tiles

EP

spanning multiple mm) for reliable linkages of processing-microstructure-property of additively manufactured TiAl, similar to earlier findings for Ti-6Al-4V [73]–[75].

AC C

Figure 12 summarizes compression true stress-strain curves taken along the build

direction for both as-deposited and HIPped material. HIPping appears to reduce the variation in compression results between samples while also reducing the flow stress. No external cracks were evident on the sides of any of the cylindrical compression samples at any strain, likely a result of the very fine grain size present in these materials and the resulting improved fracture resistance and/or smaller flaws created, particularly in compression. In addition, the average

10 | P a g e

ACCEPTED MANUSCRIPT

elastic modulus calculated from the linear portion of the compression stress-strain curves was in the range of 150-160 GPa which is in the range of cast γ TiAl [3], [5].

RI PT

Fracture toughness results, Rockwell C hardness and compression yield strength values for the present alloy are summarized in Table 2. Both the hardness and compression tests reveal that HIP reduces the strength of the AM material to values closer to that of the conventionally cast material. Increasing the notch radius from 100 µm to 500 µm increased the notch toughness,

SC

as shown for many other materials. HIP increased the average notch toughness (e.g. from 24.1

M AN U

MPa√m to 27.8 MPa√m) while reducing the degree of scatter. Both as-deposited and HIP notch toughness values were similar to that reported for the conventionally cast material tested in the TL (transverse longitudinal) orientation [68]. The current work also evaluated the fracture toughness at the midpoint of the build (i.e. midway between the start and the end of the build) as well as near start and near end of the build. Significant spatial variability in the toughness values

TE D

was observed at each of the locations with respect to the build shown in Figure 2. Representative SEM fracture surface images of the notch toughness tests for the as-

EP

deposited materials are provided in Figure 13. In general, a faceted appearance is evident with isolated regions of local plasticity in some regions, Figure 13 (a-d). Also evident are regions of

AC C

porosity that are comparable to the starting powder size. Scanning laser confocal microscopy was also conducted on the fracture surfaces in order to determine the level of fracture surface roughness. Figure 13 (e) provides a low magnification topographic map of a 1.2 mm by 1.2 mm region of the fracture surface and reveals a relatively flat fracture surface, likely a result of the very fine grain size and resulting limited tortuosity of the fracture surface. The resulting scale of fracture surface roughness is on the order of 50 µm and much less than that of as-cast material [68] where surface roughness in excess of 1 mm was measured using a similar technique. 11 | P a g e

ACCEPTED MANUSCRIPT

Figure 14 and Table 3 summarize the preliminary results of fatigue crack growth testing along with data produced by others on both AM and conventionally processed material.

RI PT

4. Discussion The fine grain size exhibited in the as-deposited EBM material is unique and, in some cases, much finer than that obtained via other commercial manufacturing techniques. Cast

SC

materials can exhibit α grain sizes in the range of 300-1000 µm [68] while wrought materials have been reported to exhibit colony sizes in the fully-lamellar condition of ~50-300 µm [76],

M AN U

[77] and γ grain sizes of ~10 µm in the equiaxed-γ condition. The as-deposited EBM material in this study exhibits an average grain size of only 21 µm and micro-hardness of 3.61 GPa for the coarsest part of the microstructure, with far smaller grain sizes and higher micro-hardness of 4.34 GPa in other regions (e.g. 3-4 µm) shown in Figure 11. Other work [38] conducted on the as-

TE D

deposited material using a similar near-γ titanium-aluminide revealed an average grain size of around 39.1 µm for the larger grain regions and 8 µm for the fine grain regions. Similarly, for HIPped material the average grain size reported for the larger grain regions was about 49.7 µm

EP

and 10.5 µm for the smaller grain region. Furthermore, the fine grains exhibited in the two-phase dendritic regions of the as-deposited EBM material exhibited a size comparable to that typically

AC C

found in wrought near-γ titanium aluminide alloys synthesized initially in ingot form [78]. While Todai et al [52] recently used optical metallography to detect the presence of a layered microstructure in TiAl made by EBM, the heterogeneous duplex microstructure shown by both large area BSE and EBSD in the present work (Figure 5 - Figure 11) is the first to be reported in the open literature for AM of near-γ TiAl in the as-deposited and HIP conditions. Similar layered microstructures have also been reported in GTAW-processed TiAl [26] that was

12 | P a g e

ACCEPTED MANUSCRIPT

conducted under different combinations of power and velocity compared to the present powder bed technique. Similar non-uniform microstructures were previously reported for cast + HIP + hot worked [66] materials, as well as those manufactured by powder metallurgy [79]. For the

RI PT

former ref. [66], the authors concluded that the coarse single-phase γ grains evolved from interdendritic regions which were the last to solidify during casting, while the two-phase bands evolved via the decomposition of a dendritic phase to produce a lamellar structure that was

SC

spherodized during hot working. The development and persistence of a similar microstructure in the PM product provided evidence that even the relatively high cooling rate imposed during gas-

M AN U

atomization synthesis is insufficient to prevent micro-segregation in powder particles which persist during subsequent HIP and hot working in the two-phase (α2 + γ) field. It worth noting that the AM material in the as-deposited condition exhibited limited lamellar features and fine equiaxed γ grains (Figure 3), as discussed below.

TE D

In the present case, the presence of a heterogeneous microstructure in the as-deposited material can also be surmised to be a result of micro-segregation [78] which persists through the fusion of powder particles during electron beam melting, solidification, and subsequent HIP in

EP

the two-phase field. Such a persistence of the microstructural heterogeneity is consistent with the

AC C

fact that true chemical homogenization can only occur if the material HIPped is raised in temperature to the single phase (α) field, which was not the case for the material in the current study. By analogy with the observations for PM γ-TiAl, it may be surmised that the microsegregation developed in the EBM process resulted from cooling/solidification rates from the liquid state that may have been comparable to those experienced during gas atomization. Two factors, however, may contribute to somewhat slower cooling rates during EBM: (1) the relatively high powder-bed preheat temperature (1050°C), which is required to minimize thermal 13 | P a g e

ACCEPTED MANUSCRIPT

stresses and thermal cracking and (2) the insulating effect of the porous, unmelted powder particles which surround the melted-and-solidified material in the powder bed. As mentioned above, the persistence of the microstructural heterogeneity throughout the build during

RI PT

deposition per se as well as subsequent HIP is consistent with the fact that chemical homogenization can only occur if the material is heated to a temperature in the single-phase (α) field for a time which is typically of the order tens of minutes or hours [54]. Thus, it is unlikely

SC

that thermal transients into the α field experienced by an already solidified layer (due to the fusion of layer(s) above it) would provide sufficient time for homogenization. Likewise, post-

M AN U

deposition HIP at 1200°C, which lies well within the α + γ filed, would certainly not lead to chemical and, hence, microstructural homogeneity. Further heat treatment evaluations are now underway in an effort to homogenous the microstructure.

Both the compression and hardness results (Figure 12 and Table 2, respectively)

TE D

exhibited reductions after HIPping. HIP-induced reductions in the strength of as-deposited AM materials are typically obtained as reviewed recently [25] for a range of materials systems.

EP

Detailed microstructural analyses on EBM-processed Ti-6Al-4V [74] have shown a coarsening of the alpha laths after HIP along with reduced strengths.

AC C

The strength/hardness of the AM material remains greater than that reported for as-cast TiAl summarized elsewhere [68] and in Table 2 that exhibited very large (e.g., mm-size) grains. While µCT of the as-deposited material revealed defects of various sizes in different locations in the build, Figure 15a, Figure 16, and Table 4, HIPping removed most of the defects as shown in Figure 15b at the currently achieved resolution. Generally, HIP induces loss of strength/hardness by microstructural coarsening, [25], [74].

14 | P a g e

ACCEPTED MANUSCRIPT

The Figure 17, taken at the same magnification, show representative regions from the asdeposited and HIPped material. It is clear that volume fraction and size of α2 phase (white phase in BSE images) has increased after the HIP treatment. The lamellar microstructure evident in the

RI PT

as-deposited material was also eliminated by the HIP treatment (due to static spheroidization) and we believe this contributes to the reduction of the strength exhibited in the HIPped material. In addition, intersecting linear features within γ grains in the as-deposited material were

SC

observed. These can be either fine twins or slip bands (marked with white line in Figure 17e). These features reduced the free slip distance within γ grains and hence may cause an increase in

M AN U

the material strength in the as-built condition. After HIP, these features were not as frequent within γ grains as it was in the as-built condition which increases the free slip distance within HIP γ grains and hence may also contribute to the reduction in the material strength. Further microstructure characterization and analysis will be conducted in the future to investigate the

TE D

mechanism of strengthening due to the intra-γ grain linear features. Thus, we believe that the drop in flow stress appears to be associated with the changes in the microstructure before and after HIP, especially in the regions with both γ and α2, Figure 3 and Figure 5.

EP

Based on the Hall-Petch equation, hypothetically a change of γ grain size from 1 to 2 µm (in the two-phase region) would lead to a change in strength of 0.3 ky (the coefficient in the

AC C

Hall-Petch equation). However, a change of γ grain size in the single-phase band from 5 to 10 µm would result in a change in strength of only 0.13 ky. In particular, in the as-built condition, the two-phase regions appear to have some fine lamellar features and fine equiaxed γ grains. Spheroidization and coarsening should both reduce the Hall-Petch contribution to strength. Due to the inherent limitations present when using 2D images to describe the 3D defect characteristics (including spatial distribution), videos for the data in Figure 15a and Figure 15b 15 | P a g e

ACCEPTED MANUSCRIPT

are provided in the supplementary material. Figure 15 and the associated 3D animation reveal that larger pores tend to spatially cluster at approximately the mid-radius region, while the core and external volume seem to have smaller pores or no pores at all. This observation matches the

RI PT

trends reported in our prior statistical analysis of µCT generated for Ti-6Al-4V also processed by EBM [74] but on larger samples. The consistency of producing core and outer regions almost free of defects while the middle radius regions have the largest defects suggest that the current

SC

and previous processing parameters and scanning strategies likely impacted the spatial distribution of defects. Further analyses of the effects of changes in processing parameters on the

M AN U

spatial distribution of thermal profiles are required to establish a correlation between the observed heterogeneous morphology and the associated spatial distribution of defects. These defects are not expected to significantly affect the compressive strength or hardness in the orientation tested, although they do appear to affect both the magnitude and scatter obtained for

TE D

the fracture critical properties tested perpendicular to the build direction as discussed later. The notch fracture toughness tests were conducted perpendicular to the build direction. Figure 18, and Table 2 show an effect of test geometry (i.e. cylindrical as-deposited vs.

EP

machined) as well as HIP on the AM material but in the range of that obtained on conventionally

AC C

cast TiAl [68]. The average results shown in Table 2 also includes the standard deviations for each condition. Higher standard deviations reflect a larger spatial variability. This spatial variability in properties is partly due to the microstructure heterogeneity shown in Figure 19. The layered microstructure shown in Figure 19 exhibits spatially varying microstructure characteristics (e.g. grain boundaries, twin boundaries, γ grain size, α2 morphology and crystallography). This heterogeneity has an impact on the location-specific mechanical properties depending on the test orientation. In general, these can include strength, ductility, and fatigue

16 | P a g e

ACCEPTED MANUSCRIPT

crack growth rate. The authors believe that such microstructure spatial heterogeneity will affect the location-specific material properties such as fracture toughness. The exact correlations are beyond the scope of the present work and the authors intend on building these correlations in

RI PT

future efforts.

The increase in notch toughness with an increase in notch radius is typically observed for metallic materials and was also shown for as-cast TiAl [68]. However, the as-deposited round

SC

samples exhibited the lowest average notch toughness and highest scatter as shown in Figure 18

M AN U

and Table 2. This likely results from a combination of the defects shown in Figure 15, Figure 16, and Table 4 in addition to including the whole as-deposited (i.e. round) sample where the higher porosity outer regions of the sample remained. Removal of the higher porosity regions in the square samples increased the notch toughness and reduced the scatter of the as-deposited materials, Figure 18 and Table 2. The highest notch toughness values and least scatter were

TE D

obtained in the square HIP samples, Figure 18 and Table 2. While the reduction in defect density and strength produced via HIP clearly increased the toughness and decreased the scatter for the notch toughness samples, the large notch radius (e.g. 100 µm, 500 µm) produces a plastic zone at

EP

fracture that effectively samples > 50 µm of the microstructure for the present toughness/strength

AC C

combinations. Thus, despite the banded nature of the microstructure shown in Figure 3 - Figure 10 and Figure 20, cracking in the notched samples will likely select the least fracture resistant path as suggested in Figure 20 and shown by the fractography in Figure 13. Fatigue precracked samples tested in fatigue and/or toughness could exhibit different trends due to the different (i.e. smaller) plastic zone sizes and sampling volumes, as discussed below. The representative fatigue crack growth data (da/dN vs. ∆K) shown in Figure 14 and Table 3 reveals location-dependent behavior, very high Paris slopes, and much lower stress 17 | P a g e

ACCEPTED MANUSCRIPT

intensity at catastrophic failure (i.e. compared to the notch toughness data). The low ductility and limited crack path tortuosity revealed by the fractography/topography, Figure 13 and Figure 21, in both the as-deposited and HIPped conditions produced rapid fatigue crack growth once the

RI PT

crack started to grow, producing the very high Paris slopes and catastrophic fracture soon after crack initiation. Figure 21 shows the generally brittle fracture surface features present on the fatigue samples as well as a lack of any fatigue striations. Closer examination of the fracture

SC

surfaces in Figure 21 reveals the bi-modal nature of the facet sizes, consistent with the heterogeneous microstructures shown in Figure 3 - Figure 10, Figure 19 and Figure 20. Very

M AN U

similar observations were reported by Patriarca [49] on a similar alloy system, but after preoxidation at 650°C for 20 hours as indicated in Table 3.

Table 5 provides estimates of the plastic zone size at different regions (e.g. threshold, Paris slope, overload) of the fatigue crack growth curve. The plastic zone sizes range from only

TE D

8.5 – 34 µm. Comparison to the levels of microstructural heterogeneity previously shown in Figure 3 - Figure 10 and Figure 20 suggests that the location-dependent fatigue crack growth properties shown in Figure 14 and Table 3 could partly result from sampling finer vs coarser

EP

grained regions in different regions of the build. Figure 19 demonstrates the level of

AC C

heterogeneity present in different regions along the length of one tall build where various properties were measured with respect to location. Similar location- and orientation-dependent fracture critical properties have been reported recently [25], [74], [80] and attributed to a complex competition between microstructure-dominated vs defect-dominated events, depending on the mechanical property under investigation.

18 | P a g e

ACCEPTED MANUSCRIPT

5. Conclusions Microstructure analysis, fracture toughness testing, and fatigue crack growth behavior of

RI PT

EBM γ-TiAl in the as-deposited and HIPped conditions were conducted in the present work. The following was found: •

Both as-deposited and HIPped AM materials exhibited heterogeneous microstructures

SC

with alternating bands of fine and coarse γ grains. These likely evolve as a result of dendritic solidification during the deposition process and the absence of sufficient

•

M AN U

subsequent exposure time in the single-phase α field for homogenization. The as-deposited material exhibited numerous isolated micro-cracks at various regions throughout the build. These were detected in metallographic sections as well as via tomography. The microstructure examination revealed that these cracks were primarily

TE D

in areas with coarse, single-phase γ grains and on many occasions within the γ grains crossing twin boundaries. Fracture surface examinations of failed as-deposited samples also revealed regions of isolated porosity, consistent with the tomography and

HIP was effective in reducing/eliminating process-induced defects but did not

AC C

•

EP

metallography results.

significantly change the grain size or reduce microstructure heterogeneities. Some residual cracks after HIP were detected and were present within the largest γ grains as confirmed by high-resolution BSE images. Microstructure homogenization will require heat treatment/HIP in the single-phase alpha regime, which unfortunately may lead to a coarse-grain, fully lamellar microstructure.

19 | P a g e

ACCEPTED MANUSCRIPT

•

Strength and hardness values of the as-deposited material exceeded that obtained on conventionally cast TiAl, consistent with the finer microstructural scale present in the AM material tested presently. HIP caused slight changes in the grain size of the single γ

regions,

elimination

of

the

lamellar

colonies

(perhaps

via

RI PT

phase

a

spheroidization/termination migration process) with coarsening of the fine γ grains in the two-phase regions. The reduced strength in the HIP material is consistent with these The presence of annealing twins was only detected within

•

M AN U

coarse γ-grains in as-deposited condition.

SC

microstructure changes.

The fracture surfaces from both notch toughness and fatigue crack growth samples showed limited plasticity and roughness along with isolated porosity in as-deposited material, with notch toughness values in the range of those reported for as-cast material.

material. •

TE D

HIP increased the notch toughness and reduced the scatter compared to the as-deposited

The preliminary fatigue crack growth tests revealed rapid crack growth after fatigue

EP

initiation that produced very high values for Paris slope and catastrophic fracture soon after fracture initiation, consistent with the limited roughness and brittle nature of the

AC C

fracture surfaces. Values for Paris slope were higher than that reported for conventionally cast material. Valid fatigue thresholds were not obtained due to difficulty in initiating fatigue cracks along with the high Paris law exponents that produced catastrophic fracture soon after fracture initiation.

•

Kmax at overload in the fatigue tests was less than that obtained in the notched AM material and less than that obtained in the conventionally cast material. Differences in

20 | P a g e

ACCEPTED MANUSCRIPT

the microstructure scale and heterogeneity throughout the build contributed to these observations, in addition to the smaller plastic zone sizes (and sampling volume) present.

RI PT

Acknowledgements: This work was partially supported by the Arthur P Armington Professorship (JJL) and conducted in the Advanced Manufacturing and Mechanical Reliability Center (AMMRC) at

SC

CWRU. The materials in this study were manufactured and provided by Arcam AB. Microstructural characterization and analytics were internally funded by Materials Resources

M AN U

LLC.

References

X. Wu, “Review of alloy and process development of TiAl alloys,” Intermetallics, vol. 14, no. 10–11, pp. 1114–1122, Oct. 2006.

[2]

Y.-W. Kim, “Intermetallic alloys based on gamma titanium aluminide,” JOM, vol. 41, no. 7, pp. 24–30, 1989.

[3]

P. Bartolotta, J. Barrett, T. Kelly, and R. Smashey, “The use of cast Ti−48Al−2Cr−2Nb in Jet Engines,” JOM, vol. 49, no. 5, pp. 48–50, 1997.

[4]

S. C. Huang and M. F. X. Gigliotti, “Chromium-Modified Titanium Aluminum Alloys and Method of Preparation.” US Patent, 1989.

[5]

F. Appel, J. D. H. Paul, and M. Oehring, Gamma Titanium Aluminide Alloys Science and Technology. Wiley-VCH Verlag, 2011.

[6]

Y. W. Kim, W. Smarsly, J. Lin, D. M. Dimiduk, and F. Appel, Eds., Gamma Titanium Aluminide Alloys. Hoboken, NJ: John Wiley & Sons, Inc, 2014.

[8]

[9]

EP

AC C

[7]

TE D

[1]

A. W. Sommer, G. C. Keijzers, Y.-W. Kim, H. Clemens, and A. H. Rosebberger, Eds., Gamma Titanium Aluminides. Warrendale, PA: TMS, 2003. V. A. C. Haanappel, H. Clemens, and M. F. Stroosnijder, “The high temperature oxidation behaviour of high and low alloyed TiAl-based intermetallics,” Intermetallics, vol. 10, no. 3, pp. 293–305, 2002.

H. Clemens, W. Smarsly, and S. Mayer, “Advanced intermetallic titanium aluminides,” in Proceedings of the 13th World Conference on Titanium, V. Venkatesh, A. Pilchak, J. Allison, S. Ankem, R. Boyer, J. Christodoulou, H. Fraser, A. Imam, Y. Kosaka, H. Rack, A. Chatterjee, and A. Woodfield, Eds. Hoboken, NJ: John Wiley & Sons, Inc., 2016, pp. 1193–1200. 21 | P a g e

ACCEPTED MANUSCRIPT

H. Clemens and S. Mayer, “Design, processing, microstructure, properties, and applications of advanced intermetallic TiAl alloys,” Adv. Eng. Mater., vol. 15, no. 4, pp. 191–215, 2013.

[11]

M. Kastenhuber, B. Rashkova, H. Clemens, and S. Mayer, “Enhancement of creep properties and microstructural stability of intermetallic β-solidifying γ-TiAl based alloys,” Intermetallics, vol. 63, pp. 19–26, 2015.

[12]

B. P. Bewlay, S. Nag, A. Suzuki, and M. J. Weimer, “TiAl alloys in commercial aircraft engines,” Mater. High Temp., vol. 3409, no. July, pp. 1–11, 2016.

[13]

H. Clemens and W. Smarsly, “Light-Weight Intermetallic Titanium Aluminides – Status of Research and Development,” Adv. Mater. Res., vol. 278, pp. 551–556, 2011.

[14]

B. F. Appel, U. Brossmann, U. Christoph, S. Eggert, P. Janschek, U. Lorenz, J. Müllauer, M. Oehring, and J. D. H. Paul, “Recent Progress in the Development of Gamma Titanium Aluminide Alloys,” Adv. Eng. Mater., vol. 2, no. 11, pp. 699–720, 2000.

[15]

M. Thomas, “Processing and Characterization of TiAl- based Alloys : Towards an Industrial Scale,” J. Aerosp. Lab, no. 3, pp. 1–11, 2011.

[16]

M. J. Weimer and T. J. Kelly, “TiAl Alloy 48Al-2Nb-2Cr Material database and application status,” in 3rd International Workshop on γ- TiAl technologies, 2006, pp. 1– 15.

[17]

B. P. Bewlay, M. Weimer, T. Kelley, A. Suzuki, and P. R. Subramanian, “The science, technology, and implementation of TiAl alloys in commercial aircraft engines,” in MRS Symp. Intermetallic-Based Alloys- Science, Technology and Applications, 2013, vol. 1516, pp. 49–58.

[18]

A. V. Fadeev, V. E. Bazhenov, and A. V. Koltygin, “Improvement in the casting technology of blades for aviation gas-turbine engines made of TNM-B1 titanium aluminide alloy produced by induction crucible melting,” Russ. J. Non-Ferrous Met., vol. 56, no. 1, pp. 26–32, 2015.

[19]

H. Saari, J. Beddoes, D. Y. Seo, and L. Zhao, “Development of directionally solidified γTiAl structures,” Intermetallics, vol. 13, no. 9, pp. 937–943, Sep. 2005.

[20]

D. J. Jarvis and D. Voss, “IMPRESS Integrated Project—An overview paper,” Mater. Sci. Eng. A, vol. 413–414, pp. 583–591, Dec. 2005.

[21]

M. Thomas, J. L. L. Raviart, and F. Popoff, “Cast and PM processing development in gamma aluminides,” Intermetallics, vol. 13, no. 9, pp. 944–951, Sep. 2005.

[23]

SC

M AN U

TE D

EP

AC C

[22]

RI PT

[10]

D. Srivastava, “Microstructural characterization of the γ-TiAl alloy samples fabricated by direct laser fabrication rapid prototype technique,” Bull. Mater. Sci., vol. 25, no. 7, pp. 619–633, 2002.

Y.-W. Kim and D. M. Dimiduk, “Progress in the understanding of gamma titanium aluminides,” JOM, vol. 43, no. 8, pp. 40–47, 1991.

[24]

B. Dogan, D. Schöneich, K.-H. Schwalbe, R. Wagner, D. Schiineich, K.-H. Schwalbe, R. Wagner, and K. Fq, “Fatigue precracking and fracture toughness testing of TiAl intermetallics,” Intermetallics, vol. 4, no. 1, pp. 61–69, 1996.

[25]

J. J. Lewandowski and M. Seifi, “Metal Additive Manufacturing: A Review of 22 | P a g e

ACCEPTED MANUSCRIPT

Mechanical Properties,” Annu. Rev. Mater. Res., vol. 46, pp. 151–186, 2016. Y. Ma, D. Cuiuri, N. Hoye, H. Li, and Z. Pan, “The effect of location on the microstructure and mechanical properties of titanium aluminides produced by additive layer manufacturing using in-situ alloying and gas tungsten arc welding,” Mater. Sci. Eng. A, vol. 631, pp. 230–240, 2015.

[27]

Y. Ma, D. Cuiuri, C. Shen, H. Li, and Z. Pan, “Effect of interpass temperature on in-situ alloying and additive manufacturing of titanium aluminides using gas tungsten arc welding,” Addit. Manuf., vol. 8, pp. 71–77, 2015.

[28]

L. Loeber, S. Biamino, U. Ackelid, S. Sabbadini, P. Epicoco, P. Fino, and J. Eckert, “Comparison of Selective Laser and Electron Beam Melted Titanium Aluminides,” in Solid Freeform Fabrication Proceedings, 2011, pp. 547–556.

[29]

L. Löber, F. P. Schimansky, U. Kühn, F. Pyczak, and J. Eckert, “Selective laser melting of a beta-solidifying TNM-B1 titanium aluminide alloy,” J. Mater. Process. Technol., vol. 214, no. 9, pp. 1852–1860, 2014.

[30]

W. Liu and J. N. Dupont, “In-situ reactive processing of nickel aluminides by laserengineered net shaping,” Metall. Mater. Trans. A, vol. 34, no. 11, pp. 2633–2641, 2003.

[31]

J. Gussone, Y.-C. Hagedorn, H. Gherekhloo, G. Kasperovich, T. Merzouk, and J. Hausmann, “Microstructure of γ-titanium aluminide processed by selected laser melting at elevated temperatures,” Intermetallics, vol. 66, pp. 133–140, 2015.

[32]

W. Li, J. Liu, S. Wen, Q. Wei, C. Yan, and Y. Shi, “Crystal orientation , crystallographic texture and phase evolution in the Ti – 45Al – 2Cr – 5Nb alloy processed by selective laser melting,” Mater. Charact., vol. 113, pp. 125–133, 2016.

[33]

L. E. Murr, S. M. Gaytan, A. Ceylan, E. Martinez, J. L. Martinez, D. H. Hernandez, B. I. Machado, D. A. Ramirez, F. Medina, S. Collins, and R. B. Wicker, “Characterization of titanium aluminide alloy components fabricated by additive manufacturing using electron beam melting,” Acta Mater., vol. 58, no. 5, pp. 1887–1894, 2010.

[34]

J. Hernandez, L. E. Murr, S. M. Gaytan, E. Martinez, F. Medina, and R. B. Wicker, “Microstructures for Two-Phase Gamma Titanium Aluminide Fabricated by Electron Beam Melting,” Metallogr. Microstruct. Anal., vol. 1, no. 1, pp. 14–27, Feb. 2012.

[35]

W. Ge, F. Lin, and C. Guo, “The effect of scan pattern on microstructure evolution and mechanical properties in electron beam melting Ti47Al2Cr2Nb,” in Solid Freeform Fabrication Proceedings, 2014, pp. 501–513.

[37]

[38]

SC

M AN U

TE D

EP

AC C

[36]

RI PT

[26]

J. Schwerdtfeger and C. Körner, “Selective electron beam melting of Ti-48Al-2Nb-2Cr: Microstructure and aluminium loss,” Intermetallics, vol. 49, pp. 29–35, 2014. S. Biamino, B. Kloden, T. Weibgarber, B. Kieback, and U. Ackelid, “Titanium aluminides for automotive applications processed by electron beam melting,” in MPIF (Metal Powder Industries Federation), 2014, pp. 96–103. S. F. Franzen, J. Karlsson, R. Dehoff, U. Ackelid, O. Rios, C. Parish, and W. Peters, “Microstructural Properties of Gamma Titanium Aluminide Manufactured by Electron Beam Melting,” in TMS Proceedings (The Minerals, Metals & Materials Society), 2011, pp. 455–462. 23 | P a g e

ACCEPTED MANUSCRIPT

S. Sabbadini, O. Tassa, P. Gennaro, and U. Ackelid, “Additive Manufacturing of Gamma Titanium Aluminide Parts By Electron Beam Melting,” in TMS Proceedings (The Minerals, Metals & Materials Society), 2010, pp. 267–274.

[40]

H. P. Tang, G. Y. Yang, W. P. Jia, W. W. He, S. L. Lu, and M. Qian, “Additive manufacturing of a high niobium-containing titanium aluminide alloy by selective electron beam melting,” Mater. Sci. Eng. A, vol. 636, pp. 103–107, 2015.

[41]

W. Ge, C. Guo, and F. Lin, “Effect of Process Parameters on Microstructure of TiAl Alloy Produced by Electron Beam Selective Melting,” Procedia Eng., vol. 81, no. October, pp. 1192–1197, 2014.

[42]

J. Porter, J. Wooten, O. Harrysson, and K. Knowlson, “Digital manufacturing of gammaTiAl by electron beam melting,” Mater. Sci. Technol. Conf., pp. 1434–1441, 2011.

[43]

T. Mclouth, Y. Chang, J. Wooten, and J. Yang, “The Effects of Electron Beam Melting on the Microstructure and Mechanical Properties of Ti-6Al-4V and Gamma-TiAl,” in Microscopy and Microanalysis, 2015, vol. 21, no. 588, pp. 1177–1178.

[44]

A. Seikh, A. Mohammad, E.-S. Sherif, and A. Al-Ahmari, “Corrosion Behavior in 3.5% NaCl Solutions of γ-TiAl Processed by Electron Beam Melting Process,” Metals (Basel)., vol. 5, no. 4, pp. 2289–2302, 2015.

[45]

S. Biamino, A. Penna, U. Ackelid, S. Sabbadini, O. Tassa, P. Fino, M. Pavese, P. Gennaro, and C. Badini, “Electron beam melting of Ti–48Al–2Cr–2Nb alloy: Microstructure and mechanical properties investigation,” Intermetallics, vol. 19, no. 6, pp. 776–781, Jun. 2011.

[46]

D. Cormier, O. Harrysson, T. Mahale, and H. West, “Freeform Fabrication of Titanium Aluminide via Electron Beam Melting Using Prealloyed and Blended Powders,” Res. Lett. Mater. Sci., vol. 2007, pp. 1–4, 2007.

[47]

M. Filippini, S. Beretta, L. Patriarca, G. Pasquero, and S. Sabbadini, “Defect tolerance of a gamma titanium aluminide alloy,” Procedia Eng., vol. 10, pp. 3677–3682, 2011.

[48]

M. Terner, S. Biamino, D. Ugues, S. Sabbadini, P. Fino, M. Pavese, and C. Badini, “Phase transitions assessment on γ-TiAl by Thermo Mechanical Analysis,” Intermetallics, vol. 37, pp. 7–10, Jun. 2013.

[49]

L. Patriarca, “Fatigue crack growth of a gamma titanium aluminide alloy,” in 9th Youth Symposium on Experimental Solid Mechanics, 2010, pp. 36–39.

[51]

[52]

SC

M AN U

TE D

EP

AC C

[50]

RI PT

[39]

M. Filippini, S. Beretta, C. Içöz, and L. Patriarca, “Effect of the Microstructure on the Fatigue Strength of a TiAl Intermetallic Alloy Produced by Additive Manufacturing,” in MRS Proceedings, 2015, vol. 1760, pp. 3–8.

M. Filippini, S. Beretta, L. Patriarca, G. Pasquero, and S. Sabbadini, “Fatigue Sensitivity to Small Defects of a Gamma–Titanium–Aluminide Alloy,” J. ASTM Int., vol. 9, no. 5, pp. 1–12, 2012.

M. Todai, T. Nakano, T. Liu, H. Y. Yasuda, K. Hagihara, K. Cho, M. Ueda, and M. Takeyama, “Effect of building direction on the microstructure and tensile properties of Ti48Al-2Cr-2Nb alloy additively manufactured by electron beam melting,” Addit. Manuf., vol. 13, pp. 61–70, 2017. 24 | P a g e

ACCEPTED MANUSCRIPT

GE, “This Electron Gun Builds Jet Engines.” [Online]. Available: http://www.gereports.com/post/94658699280/this-electron-gun-builds-jet-engines/. [Accessed: 01-Jan-2016].

[54]

K. Frkáňová and J. Lapin, “Relationship Between Microstructure and Cooling Rate in AirHardenable TiAl-Based Alloy,” in 21st International Conference on Metallurgy and Materials, 2012.

[55]

Materials Resources LLC, “MiCloud,” 2015. .

[56]

Y. Murakami, Stress intensity factors handbook. Oxford: Pergamon Press, 1987.

[57]

ASTM Standard E399, “Standard Test Method for Linear-Elastic Plane-Strain Fracture Toughness KIC of Metallic Materials,” in ASTM Book of Standards, West Conshohocken, PA: ASTM International, 2012.

[58]

ASTM Standard E647, “Standard Test Method for Measurement of Fatigue Crack Growth Rates,” in ASTM Book of Standards, West Conshohocken, PA: ASTM International, 2013.

[59]

S. Suresh, “Crack initiation in cyclic compression and its applications,” Eng. Fract. Mech., vol. 21, no. 3, pp. 453–463, 1985.

[60]

S. Suresh, L. Ewart, M. Maden, W. S. Slaughter, and M. Nguyen, “Fracture toughness measurements in ceramics: Pre-cracking in cyclic compression,” J. Mater. Sci., vol. 22, no. 4, pp. 1271–1276, 1987.

[61]

S. Suresh and J. R. Brockenbrough, “Theory and Experiments of Fracture in Cyclic Compression: Single Phase Ceramics, Transforming Ceramics and Ceramic Composites,” Acta Metall., vol. 36, no. 6, pp. 1455–1470, 1988.

[62]

G. Bergmann and H. Vehoff, “Precracking of NiAl Single Crystals by CompressionCompression Fatigue And Its Application To Fracture Tougness Testing,” Scr. Metall. Mater., vol. 30, no. 8, pp. 969–974, 1994.

[63]

J. Alcala and M. Anglada, “Fatigue and Static Crack Propagation in Yttria-Stabilized Tetragonal Zirconia Polycrystals : Crack Growth Micromechanisms and Precracking Effects,” J. Am. Ceram. Soc., vol. 80, no. 11, pp. 2759–2772, 1997.

[64]

K. J. Marsh, R. A. Smith, and R. O. Ritchie, Fatigue Crack Measurement Techniques and Applications. London, UK: Engineering Materials Advisory Services, 1991.

[65]

P. Moore, Nondestructive Testing Handbook: Volume 4 , Radiographic Testing, 3rd ed. American Society for Nondestructive Testing, 2002.

[67]

SC

M AN U

TE D

EP

AC C

[66]

RI PT

[53]

S. L. Semiatin and P. A. McQuay, “Segregation and homogenization of a near-gamma titanium aluminide,” Metall. Trans. A, vol. 23, no. 1, pp. 149–161, 1992. C. McCullough, J. J. Valencia, C. G. Levi, and R. Mehrabian, “Phase equilibria and solidification in Ti-Al alloys,” Acta Metall., vol. 37, no. 5, pp. 1321–1336, 1989.

[68]

M. S. Dahar, S. M. Seifi, B. P. Bewlay, and J. J. Lewandowski, “Effects of test orientation on fracture and fatigue crack growth behavior of third generation as-cast Ti–48Al–2Nb– 2Cr,” Intermetallics, vol. 57, no. 2, pp. 73–82, Feb. 2015.

[69]

M. Seifi, I. Ghamarian, P. Samimi, P. C. Collins, and J. J. Lewandowski, “Microstructure and Mechanical Properties of Ti-48Al-2Cr-2Nb Manufactured Via Electron Beam Melting,” in Proceedings of the 13th World Conference on Titanium, V. Venkatesh, A. 25 | P a g e

ACCEPTED MANUSCRIPT

Pilchak, J. Allison, S. Ankem, R. Boyer, J. Christodoulou, H. Fraser, A. Imam, Y. Kosaka, H. Rack, A. Chatterjee, and A. Woodfield, Eds. San Diego, CA: TMS (The Minerals, Metals & Materials Society)/Wiley, 2016, pp. 1317–1322. Materials Resources LLC Dayton OH, “MiCloud.AM.” www.MiCloud.AM. [Accessed: 28-Mar-2016].

[71]

A. A. Salem and J. B. Shaffer, “Identification and quantification of microtextured regions in materials with ordered crystal structure,” US9070203, 2013.

[72]

S. L. Semiatin, R. Nekkanti, M. K. Alam, and P. A. McQuay, “Homogenization of neargamma titanium aluminides: Analysis of kinetics and process scaleup feasibility,” Metall. Trans. A, vol. 24, no. 6, pp. 1295–1306, 1993.

[73]

M. Seifi, A. Salem, J. Beuth, O. Harrysson, and J. J. Lewandowski, “Overview of Materials Qualification Needs for Metal Additive Manufacturing,” JOM, vol. 68, no. 3, pp. 747–764, 2016.

[74]

M. Seifi, A. Salem, D. Satko, J. Shaffer, and J. J. Lewandowski, “Defect Distribution and Microstructure Heterogeneity Effects on Fracture Resistance and Fatigue Behavior of EBM Ti-6Al-4V,” Int. J. Fatigue, vol. 94, no. 1, pp. 263–287, 2017.

[75]

M. Seifi, M. Dahar, R. Aman, O. Harrysson, J. Beuth, and J. J. Lewandowski, “Evaluation of Orientation Dependence of Fracture Toughness and Fatigue Crack Propagation Behavior of As-Deposited Arcam EBM Ti-6Al-4V,” JOM, vol. 67, no. 3, pp. 597–607, 2015.

[76]

J. H. Westbrock and R. L. Fleischer, Intermetallic Compounds Principles and Practice. John Wiley & Sons, Ltd, 1995.

[77]

F. Appel, U. Lorenz, M. Oehring, U. Sparka, and R. Wagner, “Thermally activated deformation mechanisms in micro-alloyed two-phase titanium amminide alloys,” Mater. Sci. Eng. A, vol. 233, pp. 1–14, 1997.

[78]

S. L. Semiatin, G. R. Cornish, and D. Eylon, “Hot-Compression Behavior and Microstructure Evolution of Pre-Alloyed Powder Compacts of a near-Gamma Titanium Aluminide Alloy,” Mater. Sci. Eng. A, vol. 185, no. 1–2, pp. 45–53, 1994.

[79]

S. L. Semiatin, N. Frey, S. M. El-Soudani, and J. D. Bryant, “Flow softening and microstructure evolution during hot working of wrought near-gamma titanium aluminides,” Metall. Trans. A, vol. 23, no. 6, pp. 1719–1735, 1992.

Available:

SC

M AN U

TE D

EP

AC C

[80]

[Online].

RI PT

[70]

M. Seifi, M. Gorelik, J. Waller, N. Hrabe, N. Shamsaei, S. Daniewicz, and J. J. Lewandowski, “Progress Towards Metal Additive Manufacturing Standardization to Support Qualification and Certification,” JOM, vol. 69, no. 3, pp. 439–455, 2017.

26 | P a g e

ACCEPTED MANUSCRIPT

Table 1. Measured powder chemistry compared to Ti4822 alloys specification. All measurements are in wt. %. Table 2. Summary of notch toughness, hardness and compression data. Table 3. Summary of Fatigue Crack Growth results.

RI PT

Table 4. Metrics from the boxplot statistical analysis of defects’ diameter for as-deposited material showed in Figure 16. Table 5. Summary of monotonic plastic zone size (in µm) estimate for some of the test conditions at various Kmax. Plastic zone size calculation for threshold and intermediate ∆K and Kmax at failure performed based on plane strain condition due to the thickness of the sample.

SC

Figure 1. (a) Cylindrical as-built specimens, (b) machined specimens, (c) Notched toughness 3-point bend setup (crack path perpendicular to build orientation), (d) Compression specimens before and after the test (Compression axis parallel to build direction).

M AN U

Figure 2. Schematic of specimens showing the location of notch fracture toughness, compression tests, fatigue tests, and microstructural characterization experiments with respect to start and end of the build. Compression test axis was parallel to build; fracture crack path was perpendicular to build orientation. Figure 3. BSE results on the side of as-deposited specimen revealing the microstructure heterogeneity in (a) and (b). Cracks in (b) traverse multiple γ grains, (c) High resolution BSE showing the presence of lamellar colony and fine equiaxed microstructure in the duplex phase bands. Figure 4. Inverse pole figure map results on the side of HIPped specimen, (a) TiAl phase, (b) α2 phase, (c), (d) γ and α2 pole figures. Reference direction (RD) is the build direction (vertical).

TE D

Figure 5. BSE images of HIPped material revealing the drastic microstructure spatial heterogeneity under the same high magnification (a) and (b) and lower magnification (c). Various transgranular cracks are evident in (a), along with crack deflection across matrix/twin boundaries. Figure 6. (a-c) Image quality map from EBSD of the as-deposited TiAl sample showing significant spatial heterogeneity of the grain size. Build direction is bottom to top in each image.

EP

Figure 7. (a) Image quality and Build direction IPF maps of (b) coarse γ grains and (c) fine γ grains for asdeposited material shown in Figure 6(b). (d) Grain size statistics presented by boxplots shows the drastic difference in the γ grain size. Figure 8. EDS analysis of the fine vs coarse grain regions of a HIPped sample. Table of at% content shows slight difference in Al content in fine vs. coarse grain regions.

AC C

Figure 9. Image quality map of (a) As-deposited, (c) as-HIPped material showing the persistence of microstructure heterogeneity before and after HIPping process. Large area IPF map of (b) As-deposited, (d) HIPped. Figure 10. (a) Image quality and Build direction IPF maps of (b) coarse γ grains and (c) fine γ grains for HIPped material. (d) Grain size statistics presented by boxplots shows the drastic difference in the γ grain size. Figure 11. Statistical comparison between coarse and fine grained γ regions in the As-deposited and HIPped material. Figure 12. Engineering stress-strain curves showing spatial variability in σy in as-deposited materials that exhibited severe heterogeneity in the microstructure. Much less variation in flow behavior is seen after HIPping. (Compression testing was stopped at a true strain of around 0.25 as forces exceeded machine load capacity) Figure 13. (a)-(d) SEM images of fracture surface for as-deposited material after toughness experiment exhibiting various random porosities, (e) Laser confocal scanning microscopy of the fracture surface on as-deposited material showing minimal surface roughness.

27 | P a g e

ACCEPTED MANUSCRIPT

Figure 14. Fatigue crack growth behavior in the as-deposited material at two different locations and near the Middle of the build after HIPping. Figure 15. Low and high magnification views of µCT scans of SL (vertical) build showing defects for (a) asdeposited (b) HIPped specimen. The dots are measured defects. No defect was observed after HIPping.

RI PT

Figure 16. Statistical analysis using boxplot for the defect diameters measured by µCT for an as-deposited SL (vertical) sample at various locations with respect to the build showing significant spatial variation in defect population for the as-deposited specimen. No defects were detected after HIPping. Figure 17. BSE images of as-deposited and HIPped material revealing the drastic microstructure spatial heterogeneity before and after HIP treatment under the same magnification (a) and (b) and low magnification (c) and (d) high magnification (e) higher magnification exhibiting fine twins in the γ grains. The areas within the marked “very thin twins” are too small to be indexed by EBSD.

SC

Figure 18. Variation of notched fracture toughness for as-deposited and HIPped specimens indicating geometry effect as well as HIP effect. Average fracture toughness increased after HIPping with much lower variation.

M AN U

Figure 19. Image quality map after HIPping showing heterogeneity over a wider region. The scale bar on all figures is the same (240 µm). Figure 20. Crack growth is shown with red arrows in fine vs. coarse grain region. Higher fracture toughness obtained for fine grain region.

AC C

EP

TE D

Figure 21. (a)-(d) Fracture surface SEM images of HIPped sample tested to failure in fatigue. No evidence of porosity and some local ductility is evident while fracture surface features are generally brittle.

28 | P a g e

ACCEPTED MANUSCRIPT

Table 1. Measured powder chemistry compared to Ti4822 alloys specification. All measurements are in wt. %. Powder X

Aluminum

32.0 - 33.5

33.8

Chromium

2.2 - 2.6

2.3

Niobium

4.5 - 5.1 Max. 0.05

Carbon

Max. 0.025

Oxygen

Max. 0.12

Nitrogen

Max. 0.02

Hydrogen

Max. 0.003

4.7

0.02

0.01

0.07

M AN U

Iron

RI PT

Ti-48Al-2Cr-2Nb Alloy specification

SC

Element

Bal.

AC C

EP

TE D

Titanium

0.004

<0.001 Bal.

ACCEPTED MANUSCRIPT

Table 2. Summary of notch toughness, hardness and compression data. Specimen Condition

Notch Root Radius (µm)

Thickness, B (mm)

Notched Toughness Kq (MPa√m)

As-built (Round)

500

15

26.7 ± 9.6

As-built (Round)

100

15

18.6 ± 3.9

As-built (Square)

100

10

24.1 ± 6.5

As-HIPped (Square) As Cast (LT)[67]

100

10

27.8 ± 0.4

100

7.5

25

As Cast (TL) [67]

100

9.5

Rockwell C

Compression Yield Strength (MPa)

RI PT

NA 30±1

23±1

572±3

24±2

N/A

SC 24

M AN U TE D EP AC C

740±14

ACCEPTED MANUSCRIPT

Table 3. Summary of Fatigue Crack Growth results. Specimen Condition

Load Ratio, R

Threshold, (MPa√m)

∆Kth

Paris Slope, m

Kmax, (MPa√m)

0.3

8*

42

13.2

As- deposited (Near END of build)

0.3

9*

105

13.9

As-HIPped

0.1

13

Pre-oxidized [47]

0.05

6

Pre-oxidized [47]

0.6

3

As-Cast (LT) [63]

0.3

8

As-Cast (TL) [63]

0.3

9

RI PT

As- deposited (Near START of build)

18.2

15

10.8

40

11.5

23

18.3

SC

M AN U TE D EP AC C

50

34

20.4

ACCEPTED MANUSCRIPT

Size (µm)

Minimum

0.14

Lower quartile

0.21

Median

0.25

Upper quartile

0.31

Maximum

0.70

AC C

EP

TE D

M AN U

SC

Statistic

RI PT

Table 4. Metrics from the boxplot statistical analysis of defects’ diameter for as-deposited material showed in Figure 16.

ACCEPTED MANUSCRIPT

Table 5. Summary of monotonic plastic zone size (in µm) estimate for some of the test conditions at various Kmax. Plastic zone size calculation for threshold and intermediate ∆K and Kmax at failure performed based on plane strain condition due to thickness of sample.

Kmax, threshold

0.1 0.3

13 8.5

Kmax, intermediate

Plane Strain rp,

Kmax, failure

(MPa√m)

(µm)

(MPa√m)

Plane Strain rp, (µm)

16 13

15 11

21 21

18 14

31 34

AC C

EP

TE D

M AN U

SC

(MPa√m)

Plane Strain rp, (µm)

RI PT

Load Ratio (R)

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

AC C

EP

TE D

M AN U

SC

RI PT

ACCEPTED MANUSCRIPT

ACCEPTED MANUSCRIPT

Highlights: - Bulk TiAl produced via EBM

heterogeneity

RI PT

- Microstructure Characterization using large area EBSD to characterize microstructure

- Tomography characterization of as deposited defects and effects of HIP - Location dependence of fracture and fatigue behavior

AC C

EP

TE D

M AN U

SC

- Effect of HIP treatments on microstructure, defects, and resulting fracture/fatigue behavior