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ALUMINUM CONTENT AND POSTWELDED HEAT TREATMENT EFFECTS ON THE MICROSTRUCTURE EVOLUTION AND MECHANICAL PROPERTIES OF ELECTRON

BEAM WELDED CAST GAMMA TITANIUM ALUMINIDES – Part I: Microstructure Evolution of Electron Beam Welded Cast Gamma Ti-48Al-2Nb-2Cr

C.M. Jensen, H. Zhang, W.A. Baeslack III and T.J. Kelly

Abstract

Gamma TiAl alloys are of considerable interest as potential material for aerospace turbine engine applications. The ability to join these alloys usually using conventional welding techniques, such as, electron-beam(EB) method, is of great importance. In this investigation, five EB weldments, containing Al range from 46.3 to 48.3 at% produced with two heats, were examined at conditions of as-stress-relieved(SR) and post-welded heat treated(PWHT) at temperatures of 1050, 1150 and 1250°C. The microstructures of SR and PWHT of EB weldment were characterized in detail using Light(OP), Scanning Electron(SEM) and Transmission Electron(TEM) Microcopy. All EB welds exhibited a predominantly lamellar γ/α2 fusion zone(FZ) microstructure in the SR condition, with increasing proportions of equiaxed γ grains, and a decreasing volume proportion of intergranular α2 phase, observed with increasing Al content. Although PWHT did not affect base metal(BS) microstructure significantly, it promoted further recrystallization and growth of the γ grains at FZ. Increasing the PWHT temperature resulted in an increase in the volume fraction of α2 phase in the FZ. At the highest temperature, 1250°C, the presence of B2 phase was detected in FZ containing 47.1 at% Al or less. 1. INTRODUCTION

Gamma TiAl alloys, which have high ratio of strength over weight, have been considered as the most promising material for aerospace applications1-3. The effective utilization of these alloys will commonly require their joining. Unfortunately, the rapid thermal cycles experienced during fusion welding can promote the formation of weld microstructures that are metastable, and that was found to degrad room temperature ductility relative to the base metal4-8. Postweld heat treatment offers a potential for modifying the SR microstructure and thereby producing a stable microstructure with enhanced room temperature ductility. The purpose of this investigation was to examine and thereby understand the influence of the postweld heat treatment and Al contents on the resulting weld FZ microstructures, which will be reported this paper, and mechanical properties and fracture behavior of Ti-48Al-2Cr-2Nb(Ti-48-2-2) (at.%) cast gamma titanium aluminides, which will be reported on the following paper.

2. EXPERIMENTAL PROCEDURES 2.1 Base Materials

Investment cast Ti-48Al-2Cr-2Nb gamma titanium aluminide material used in this study was produced by Precision Castparts Corporation (PCC) and Howmet to GE’s specification for the “Ti-48-2-2” alloy. The actual compositions of the base metals(BS) are listed in Table 1. The BSs were heat treated at 1095°C for 4 hours, hot isostatic pressed (HIP’ed) at 1185°C for 4 hours, and a final heat treatment of 1205°C for 2 hours to form duplex structure before the welding. Table 1 Compositions (at.%) of base materials.

Weld # Ti Al Cr Nb Fe O Producer A Bal. 46.3 2.01 1.94 0.01 0.19 Howmet B Bal. 46.3 1.99 1.96 0.02 0.16 Howmet

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C Bal. 47.1 1.99 1.97 0.02 0.15 Howmet D Bal. 47.8 1.95 1.97 0.02 0.18 Howmet E Bal. 48.3 1.92 1.98 0.02 0.18 PCC

2.2 Welding Procedures

The welding parameters were developed by GEAE. The EB weldments were produced in a specially constructed vacuum chamber which allowed the plates to be heated prior to, during and following welding. Two pieces 1.5-inch by 10-inch with a ¼-inch thickness were joined along the long edge using a square butt-joint to produce autogenous, full-penetration welds. Run-on and run-off tabs, in conjunction with backing bars, were utilized to contain the molten metal and prevent it from falling through the backside of the joint. Table 2 Electron beam welding parameters.

2.3 Postweld Heat Treatments

Postweld heat treatment was performed by Wright-Patterson Air Force Base (WPAFB) in a vacuum furnace for 5 hours at the temperatures of 1050, 1150 and 1250 °C. Following heat treatment, the samples were furnace cooled to room temperature.

2.4 Characterization Techniques

The weld plate was sectioned into about 8 mm thick slices using an electric-discharge machine. The one was for SR microstructure evolution. The others were then PWHT’ed in a vacuum furnace from 1050 to 1250°C for 5 hours followed by furnace cooling. Both the SR and PWHT’ed slices were then mounted in conductive bakelite and metallographicly prepared for OP and SEM examination. The mounted samples were then examined at 20KeV in a Phillips XL-30 Scanning Electron Microscope in the as-polished condition. Samples were also optically examined in the etched condition using a 25ml H2O + 25ml HNO3 + 2ml HF + 50ml Glycerol solution.

TEM foils were removed from the specimens by electro-discharge machining cylinders from the FZ and BS, and diamond sawing thin slices from them. The slices were then ground down to less than 125 µm, followed by twin-jet polishing in a solution of 6-8 ml HClO4 + 250 ml methanol + 150 ml 2-butoxy ethanol, at a voltage of 20 V, a current of 6-8 mA, and temperature range from -30 to -40°C. TEM analysis was performed using a Philips CM 200 Analytical Transmission Electron Microscope.

3. RESULTS and CONCLUSIONS 3.1 Base Material 3.1.1 Microstructure Characterization

Figures 1(a-e) are polarized light micrographs of the base materials used in this study. Since the base materials covered a range of Al contents, from 46.3 to 48.3 %, a corresponding range of microstructures was observed. A transition from a primarily lamellar structure, in the 46.3 % Al alloy, to a near-gamma structure with a small volume fraction of lamellar colonies in the 48.3 % Al alloy was observed to occur over the range of Al contents. This observation is consistent with what is predicted from the binary phase diagram, with lower Al contents favoring higher fractions of α phase at the same heat treatment temperature used, and correspondingly higher fractions of lamellar

Current (mA) Voltage (kV) Trv. Rate (ipm) Preheat (°C) S.R. (°C)A 22.6 150 15 540 820/2hrs

B-E 22.6 150 15 650 980/2hrs

Process Parameter Thermal History

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colonies at room temperature. SEM/BSE analysis confirmed the observations made with optical microscopy while

revealing additional microstructural information. Figures 2(a-e) show the same trend of decreasing amounts of lamellar colonies and increasing proportions of γ grains, with increases in Al. Figures 2 (a) and (b), 46.3 % Al, reveal a larger proportion of angular particles located at colony and γ grain boundaries compared with others. Further examination of the high and low Al base materials, Figures 3 (a) and (b), coupled with EDS analysis, suggested the presence of B2 particles in the low Al material since these particles contain more 15 at% Cr, as indicated by the arrows in the figure. The presence of B2 phase was not detected in the base materials with Al contents at or above 47.8 Al. TEM analysis of the base materials confirmed the SEM/BSE observations, as shown in Figures 4(a-c). B2 phase was detected in the low Al content base material, located primarily at grain boundaries, Figure 4 (a). 3.2 Fusion Zones 3.2.1 As-Stress-Relieved

Five full-penetration EB welds were produced with four different Al levels and two different heats, of Ti-48-2-2. There was no obvious difference in microstructure between the weldments produced at HSR and LSR containing 46.3 at% Al using OP and SEM, so that the examination was focused on the four weldments, which produced under the same conditions with different Al contents from 46.3 to 48.3 at%.

Representative macrographs, Figures 5(a-d), of the four EB welds, produced under the same conditions with different Al contents from 46.3 to 48.3 at%, show characteristic features of a narrow weld bead with deep penetration. EB welds were produced to determine the effects of the Al content on the microstructure evolution. Even at low magnification, it is apparent that as the Al content increases, there is a corresponding increase in the volume fraction of γ grains in the FZ, as shown in Figure 6(a-d).

Detailed analysis of the EB welds revealed distinctive differences between them. Microstructural differences will be described in terms of solidification morphology, the fraction and size of γ grains, the size and proportion of lamellar colonies, and the size and distribution of α2 and B2 phases.

Polarized light microscopy is used to reveal the γ grain structure and to some extent the lamellar colonies. In Figures 6(a-d), polarized light micrographs show the differences in the relative amount of γ grains produced for the different compositions. The weld with 46.3%Al is comprised primarily of lamellar colonies, Figure 6(a), with a small volume fraction of γ grains. As the Al content increases, the fraction of γ grains tends to increase, with the welds with 47.8 and 48.3%Al, Figures 6(c) and (d), containing primarily equiaxed γ grains, with little lamellar constituent. The coarseness of the microstructure remains fairly constant in each of the welds.

Examination of the FZ’s with SEM/BSE revealed differences in the chemical composition of individual phase. Figures 7-10 show the solidification morphologies of the EB welds produced with different Al contents. A change in solidification morphology between the Al contents of 46.3 and 47.1%Al was observed, with the remaining welds showing little difference. The change observed between the 46.3 and the 47.1%Al is believed to be a transition in the solidification path from primary β to a mixed (β + α) solidification mode. From those Figures, it is evident that coarse dendrites are produced in the welds with 46.3%Al, where as those in higher Al containing material appear cellular-dendritic. This change is either due to changes in composition, affecting the solidification morphology, or changes in the solidification path which affect the morphology. Increasing Al levels will favor a transition from β to α solidification, which is believed to be the present case.

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Utilizing the updated binary phase diagram,[10] the primary path of solidification is predicted as β. It should be noted that the binary phase diagram is only a predictive tool for what actually occurs in a quaternary alloy system. However, in this study cubic dendrites were observed for all of the alloys, which indicates that solidification was indeed primary β. The alloy with the lowest Al content is believed to have undergone complete β solidification due to the observation of a β decomposition product as shown in Figure 7. The deviation from the predicted solidification paths may be as a result of two different factors. First, it has been established that alloy additions of chromium and niobium will shift solidification towards β.[11] An investigation by Naka et al determined that a combined total of chromium and niobium greater than 5 at. % will stabilize β solidification.[16] Based on these observation, the Ti-48Al-2Cr-2Nb alloy will undergo a mixed mode solidification path of β and α. The second reason for a shift in the solidification path is the loss of Al during EB welding. Aluminum has a low vapor pressure that could lead to possible loss of Al during the welding in a high vacuum. The loss of Al would shift the solidification path further towards β solidification. Alloys containing between 47.1 and 48.3 % Al exhibited similar solidification behavior to each other, having cored cubic dendrites. The cored appearance has been attributed to the peritectic reaction.[12] Due to the non-equilibrium nature of weld solidification, the solidification path will shift towards α solidification due to the sluggish nature of the peritectic reaction. These results were consistent with Lee’s work in which the solidification behavior of laser beam welds was characterized in Ti-48Al-2Cr-2Nb alloys with Al ranging from 45.5 to 47.9 %, all alloys exhibited either β solidification or a mixed β/α solidification path.[4]

The sequence of phase transformations in the EB weld FZ’s was dependent on the alloy composition and cooling rates. In alloys with 46.3 % Al content, it was suggested that they underwent β solidification. Upon further cooling, a transformation from β to α occurs. This transformation is believed to take place as a diffusional growth of α plates into β as was described to occur in the GTA weld by authors. Following the transformation to α, a lamellar γ/α2 evolved in the manor described for the GTA weld. Alloys containing greater than 47.1 % Al solidified as mixed β/α solidification. The following sequence of transformations produced a lamellar γ/α2 constituent.

Due to the application of a stress-relief heat treatment immediately following EB welding, it can only be speculated as to the exact microstructure immediately following welding. However, observations after laser beam welding by Lee have shown that in the same alloy, a fully lamellar microstructure is produced with sufficiently low cooling rates (~150°C/s).[4] Due to the preheat that was used in this investigation, it can be speculated that the cooling rates were such that a lamellar constituent would be produced. The stress-relief applied to the welds following welding is believed to cause partial recrystallization of the lamellar microstructure, which accounts for the observations that were made.

TEM characterization using dark field image techniques revealed differences in quantity of α2 phase in the lamellar colonies of the two welds produced with different preheat and stress-relief temperatures, LSR and HSR. Figure 11 shows bright-field and dark-field images of the lamellar constituent in the FZ’s. The dark-field is imaged using α2 variant such that the α2 plates are illuminated. A greater quantity of α2 phase was observed in the FZ than the BS, which was attributed to the higher cooling rates through the α phase field during welding. The high cooling rate acts to suppress the transformation of α→γ, which results in the retention of more α phase. TEM characterization of the remaining FZ’s revealed both colonies of lamellar γ/α2 and equiaxed γ grains with α2 particles. Figures 12(a-f) show representative micrographs of these regions in the 46.3, 47.1 and 48.3%Al welds. Differences in the size and morphology of the α2 particles were observed and the representative micrographs show an increase in the size of the α2 particles and an increase in the lamellar spacing for an increase in Al content.

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3.2.2 Postweld Heat-Treated Conditions Postweld heat treatment of the EB welds transforms the predominantly lamellar FZ to an

equiaxed γ grain structure with α2 particles located both inter- and intra-granularly. PWHT temperature influences the resultant γ grain size as well as the α2 particle size, distribution and morphology. For simplicity, only the PWHT temperatures of 1050 and 1250°C will be discussed, excluding the 1150°C which resulted in intermediate microstructures.

PWHT at 1050°C/5 hours resulted in the transformation of the lamellar FZ to an equiaxed γ grain structure for all Al contents examined. Figures 13(a-d) are polarized light micrographs of the FZ microstructures. In the welds produced with the two different stress-relief conditions, Figures 13(a) and (b), there is little apparent difference between the two microstructures, each exhibiting a large distribution of γ grain sizes, which tend to occur in bands. It is believed that the variation in grain size is related to the solidification substructure of the FZ. The welds produced in higher Al containing material, Figures 13(c) and (d), exhibited a more regular distribution of grain size and remained finer than the previously discussed welds. In general, no apparent trend was observed to occur with increases in Al.

Examination of welds PWHTed at 1050°C/5 hours with SEM/BSE revealed the α2 phase in the FZ microstructures. Figures 14(a-d) are high magnification micrographs which illustrate the transformation behavior of the FZ after PWHT. It is evident that recrystallization of the FZ occurred and α2 particles formed both intra- and inter-granularly. Minimal differences exist for the different stress-relife conditions, Figures 14(a) and (b), however, a decrease in the amount of α2 is observed with increases in Al content, Figures 14(b-d).

PWHT at 1250°C/5 hours produced appreciable differences in the FZ microstructure, especially at the lower Al levels. Examination of the two different stress-relief conditions, in welds produced with 46.3 %Al, Figures 15(a) and (b), showed large γ grains, greater than 100 microns, with many fine γ grains and particles throughout the microstructure. As in the previous heat treatment, no appreciable differences between the two stress-relief conditions were observed. Welds produced in higher Al contents produced a γ grain microstructure with little variability in grain sizes. Figures 15(c) and (d) are polarized light micrographs of welds produced in 47.1 and 48.3%Al titanium aluminide with little difference between the two microstructures. Grain sizes were on the order of 30 to 50 microns in diameter.

SEM/BSE analysis of the FZ’s revealed the amount and morphology of the α2 phase. Weld FZ’s in plate with 46.3%Al had a high volume fraction of α2, which was observed in the form of blocky grain boundary particles and intragranular Widmanst¬tten plates, Figures 16(a) and (b). The large amount of α2 is a result of the low Al content, favoring a higher fraction of the phase. The Widmanst¬tten plates are typically characterized as forming on-heating with an orientation relation to the γ grain. FZ microstructures with decreasing Al contents exhibited a corresponding decrease in the amount of α2 phase, as observed in Figures 16(c) and (d). The α2 morphology in the higher Al welds is predominately intergranular in nature, forming large blocky particles at γ grain boundaries.

In general, PWHT had little effect on the base metal, with the exception of the highest heat treatment temperature of 1250°C. The starting base metal microstructures in all Al contents was a duplex structure with varying fractions of lamellar colonies, depending on the Al content, as discussed previously. Figures 17(a-d) demonstrate the effect of the PHWT at 1250°C/5 hours for the different base materials investigated. Base metals with 46.3%Al exhibited significant increases in the volume fraction of α2 in both the forms of blocky grain boundary particles and Widmanstatten plates, Figures 17(a) and (b). In materials with higher Al contents, the effect was not as pronounced, however, an increase was evident in the form of intragranular Widmanst¬tten plates, Figures 17(c) and (d).

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TEM characterization of the PWHT’ed FZ’s revealed the presence of both B2 and α2 phases in the equiaxed γ grain microstructure. Figures 18(a-d) show both the low (46.3 %) and high (48.3 %) Al contents after PWHT at 1050 and 1250°C. After PWHT at 1050°C, small α2 particles were observed, primarily at γ grain boundaries as shown in Figures 18(a) and (c). PWHT at 1250°C promoted the formation of B2 phase in the low Al containing material. Figure 18(b) shows the large B2 particles adjacent to α2 plates in a γ grain matrix. The presence of B2 phase was not detected in material with Al contents at or above 47.8 %. PWHT at 1250°C promoted a coarsening of the α2 particle and increased the volume fraction.

4. Conclusions

Welding procedures have been developed by GEAE and utilized in the present study for which crack-free EB welds were successfully produced in cast Ti-48Al-2Nb-2Cr with a range of aluminum contents. The application of weld preheat and immediate stress-relief heat treatment, and a maintenance of a continuous vacuum environment, prevents cracking by minimizing thermally-induced stresses and precluding environmental affects.

The EB welded FZ’s experienced dendritic solidification for all of the Al contents investigated. A transition in solidification behavior was observed between Al contents of 46.3 and 47.1 % Al (atomic %). The transition is believed to be associated with a transition from primary β solidification in the plate containing 46.3 % Al, to a mixed β/α solidification mode in the plates containing 47.1 % Al and higher.

The stress-relieved FZ’s were comprised primarily of duplex microstructures, which were appreciably finer than the parent material. • Increasing the preheat and stress-relief temperature of the weld caused an increase in the

volume fraction of γ grains in the FZ. • Increasing aluminum content of the base plate was found to increase the volume fraction of γ

grains in the weld FZ. PWHT of the EB FZ’s promoted a transformation of the duplex microstructure to an

equiaxed γ grain microstructure. • Increases in the preheat and stress-relief temperature had no apparent effect on the resultant

microstructure after PWHT. • Increasing aluminum content resulted in a decrease in the volume fraction of α2 phase in the

FZ’s. • Increasing the PWHT temperature for the EB FZ’s resulted in an increase in the volume

fraction of α2 phase. At the highest temperature, 1250°C, the presence of B2 phase was detected in FZ’s containing 47.1 % Al or less.

References 1. C.M. Austin and T.J. Kelly, “Progress in Implementation of Cast Gamma Titanium Aluminide”,

Proceedings of the Conference Gamma Titanium Aluminides, TMS, 1995, pp. 21-32. 2. J.C. Williams, “Intermetallics for Structural Applications; Potential, Reality and the Road

Ahead”, Proceedings of the Conference International Symposium of Structural Intermetallics, TMS, 1997, pp. 3-7.

3. T.J. Kelly and W.A. Baeslack III, “Advanced Titanium-Base Material Joining Technology”, Report of the Naval Air Warfare Center, 1996.

4. J. Lee, “A Study of Weld Fusion Zone Phenomena in Cast Gamma Titanium Aluminides” Ph.D. Dissertation, The Ohio State University, 1995.

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5. L.C. Mallory, W.A. Baeslack III, D. Phillips and T.J. Kelly, Gas Tungsten-Arc Welding of a Ti-48Al-2Cr-2Nb Gamma Titanium Aluminide”, Proceedings of the Conference Titanium ’92 Science and Technology, 1993, pp. 1115-1122.

6. L.C. Mallory, W.A. Baeslack III and D. Phillips, “ Evolution of the Weld Heat-Affected Zone Microstructure in a Ti-48Al-2Cr-2Nb Gamma Titanium Aluminide”, Journal of Mat. Sci. Letters, Vol. 13, 1994, pp. 1061-1065.

7. L.C. Mallory, “A Study of Weld Heat-Affected Zone Phenomena in Cast Gamma Titanium Aluminides”, Ph.D. Dissertation, The Ohio State University, 1996.

8. P.L. Threadgill, “The Prospects for Joining Titanium Aluminides”, Mat. Sci. and Engr., A192/193, 1995, pp. 640-646.

9. Y.-W. Kim, “Intermetallic Alloys Based on Gamma Titanium Aluminide”, Journal of Materials, July 1989, pp. 24-30.

10. C. McCullough, J.J. Valencia, C.G. Levi and R. Mehrabian, “Phase Equilibrium and solidification in Ti-Al Alloys”, Acta Metall., Vol. 37, No. 5, 1989, pp. 1321-1336.

11. S. Naka, M. Thomas, C. Sanchez and T. Khan, “Development fo Third Generation Castable Gamma Titanium Aluminides: Role of Solidification Paths”, Proceedings of the Conference International Symposium of Structural Intermetallics, TMS, 1997, pp. 313-322.

12. H.D. Solomon, “Fundamentals of Weld Solidification”, ASM Handbook Volume 6, ASM International, Metals Park, Ohio, 1994, pp. 45-54.

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Figure 1 Polarized light micrographs of the EB weld base material: (a) 46.3%Al, LSR; (b) 46.3%Al, HSR; (c) 47.1%Al; (d) 47.8 Al and (e) 48.3%Al.

250 µm

(e)

250 µm

(a)

250 µm

(b)

250 µm

(c)

250 µm

(d)

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Figure 2 SEM/BSE micrographs of the EB weld base material: (a) 46.3%Al, LSR; (b) 46.3%Al, HSR; (c) 47.1%Al; (d) 47.8%Al and (e) 48.3%Al.

100 µm

(e)

100 µm

(d)

100 µm

(c)

100 µm

(b)

100 µm

(a)

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Figure 3 SEM/BSE analysis of the particles in the EB weld base materials: (a) 46.3%Al, LSR and

(b) 48.3%Al. Arrows in micrograph indicate B2 particles.

25 µm

(a)

25 µm

(b)

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Figure 4 TEM micrographs of the EB weld base materials: (a) 46.3 % Al, note B2 particle at colony/grain boundary; (b) 47.1 % Al and (c) 48.3 % Al base material.

5 µm

(a)B2

1 µm

(b)

2 µm

(c)

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Figure 5 Macrographs of the EB welds in the as-welded plus stress-relieved condition: (a) 46.3%Al, LSR; (b) 46.3%Al, HSR; (c) 47.1%Al and (d) 48.3%Al.

2.5 mm

(a)

2.5 mm

(b)

2.5 mm

(c)

2.5 mm

(d)

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Figure 6 Polarized light micrographs of the EB weld fusion zone in the as-welded plus stress-relieved condition: (a) 46.3%Al, LSR; (b) 46.3%Al, HSR; (c) 47.1%Al; (d) 47.8%Al and (e)

48.3%Al.

200 µm

(a)

200 µm

(b)

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Figure 6 Polarized light micrographs of the EB weld fusion zone in the as-welded plus stress-relieved condition: (a) 46.3%Al, LSR; (b) 46.3%Al, HSR; (c) 47.1%Al; (d) 47.8%Al and (e)

48.3%Al.

200 µm

(c)

200 µm

(d)

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Figure 6 Polarized light micrographs of the EB weld fusion zone in the as-welded plus stress-relieved condition: (a) 46.3%Al, LSR; (b) 46.3%Al, HSR; (c) 47.1%Al; (d) 47.8%Al and (e)

48.3%Al.

200 µm

(e)

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Figure 7 SEM/BSE micrographs of the high temperature stress-relief EB weld fusion zone in the 46.3%Al titanium aluminide: (a) low magnification of the solidification structure and (b) high

magnification micrograph showing the substructure in greater detail.

100 µm

(a)

20 µm

(b)

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Figure 8 SEM/BSE micrographs of the EB weld fusion zone in the 47.1%Al titanium aluminide: (a) low magnification of the solidification structure and (b) high magnification micrograph showing the

substructure in greater detail.

20 µm

(b)

100 µm

(a)

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Figure 9 SEM/BSE micrographs of the EB weld fusion zone in the 47.8%Al titanium aluminide: (a) low magnification of the solidification structure and (b) high magnification micrograph showing the substructure in greater detail.Figure 10 SEM/BSE micrographs of the EB weld fusion zone in the 48.3%Al titanium aluminide: (a) low magnification of the solidification structure; (b) and (c) high

magnification micrographs showing the substructure in greater detail.

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Figure 11 TEM bright-field and dark-field micrographs of the lamellar as-welded EB weld fusion zone microstructure regions in the: (a) and (b) LSR; and (c) and (d) HSR conditions in 46.3%Al base material.

100 nm

(a)

100 nm

(b)

100 nm

(c)

100 nm

(d)

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Figure 12 TEM micrographs illustrating the effect of aluminum content on the EB weld fusion zone microstructures in: (a) and (b) 46.3 % Al; (c) and (d) 47.1 % Al; and (e) and (f) 48.3 % Al.

200 nm

(b)

α2

200 nm

(a)

500 nm

(c)

500 nm

(d)

α2

2 µm

(e)

1 µm

(f)

α2