Elevated-Temperature Mechanical Behavior of As-Cast and Wrought Ti-6Al-4V-1B

Transcription

1 Elevated-Temperature Mechanical Behavior of As-Cast and Wrought Ti-6Al-4V-1B W. CHEN, C.J. BOEHLERT, J.Y. HOWE, and E.A. PAYZANT This work studied the effect of processing on the elevated-temperature [728 K (455 C)] fatigue deformation behavior of Ti-6Al-4V-1B for maximum applied stresses between 300 to 700 MPa (R = 0.1, 5 Hz). The alloy was evaluated in the as-cast form as well as in three wrought forms: cast-and-extruded, powder metallurgy (PM) rolled, and PM extruded. Processing caused significant differences in the microstructure, which in turn impacted the fatigue properties. The PM-extruded material exhibited a fine equiaxed a + b microstructure and the greatest fatigue resistance among all the studied materials. The b-phase field extrusion followed by cooling resulted in a strong a-phase texture in which the basal plane was predominately oriented perpendicular to the extrusion axis. The TiB whiskers were also aligned in the extrusion direction. The a-phase texture in the extrusions resulted in tensile-strength anisotropy. The tensile strength in the transverse orientation was lower than that in the longitudinal orientation, but the strength in the transverse orientation remained greater than that for the as-cast Ti-6Al-4V. The ratcheting behavior during fatigue is also discussed. DOI: /s y Ó The Minerals, Metals & Materials Society and ASM International 2011 I. INTRODUCTION W. CHEN, formerly PhD Student, Department of Chemical Engineering and Materials Science, Michigan State University, East Lansing, MI 48824, is Postdoctoral Fellow, Materials Science and Technology Division, Oak Ridge National Laboratory, Oak Ridge, TN Contact chenwei@ornl.gov C.J. BOEHLERT, Associate Professor, is with the Department of Chemical Engineering and Materials Science, Michigan State University. J.Y. HOWE and E.A. PAYZANT, Staff Members, are with the Materials Science and Technology Division, Oak Ridge National Laboratory. Manuscript submitted July 21, Article published online February 4, 2011 IN recent years, a considerable amount of research has been directed toward developing boron-modified titanium (Ti) alloys. [1 21] There are significant benefits for adding boron (B) to Ti alloys. The addition of 1 wt pct of B increases the Young s modulus (E), yield strength (YS), and ultimate tensile strength (UTS) values of Ti-6Al-4V (Ti-64) by over 15 pct both at room temperature (RT) [2] and 728 K (455 C). [3] Small B additions do not result in large density changes; thus, the higher strength and stiffness of B-modified alloys provides important increases in specific strength and stiffness. [6 8] The significant increase in strength and E values arises from the strong and stiff TiB phase that precipitates in situ during solidification. [4 6] The TiB phase effectively pins grain boundaries during the solidification process so that a fine grain structure is obtained after cooling through the a + b phase field (0.1 wt pct of B addition reduced the as-cast grain size of Ti-64 from 1700 to 200 lm). [5] This grain refinement also contributes to strengthening. Due to the similarity of the thermal expansion coefficients ( / C for Ti and / C for TiB [7] ), processing of such alloys is less challenging than for other ceramicreinforced alloys. These two features lead to significant benefits for the thermomechanical processing of B-modified Ti-64. B-modified Ti alloys can be produced via traditional ingot metallurgy (IM) and powder metallurgy (PM) processes. They can also be themomechanically processed using rolling, extruding, and forging. Toyota, Inc. (Aichi, Japan) has commercialized B-modified Ti alloys for intake and exhaust valves of automotive engines using a PM process. [8] However, in order to extend B-modified Ti alloys for elevated-temperature fracture-critical applications (such as the jet engine blades), understanding the mechanisms of deformation and fracture at elevated temperature is vital. Previous work evaluated the effect of different B additions (0, 0.1, and 1 wt pct) on the elevated-temperature fatigue behavior of as-cast and cast-and-extruded Ti-64 alloys. [3,9] It was found that B additions did not degrade the fatigue resistance of the as-cast Ti-64 alloy. [9] The cast-and-extruded Ti-64-xB alloys were significantly stronger and exhibited longer average fatigue life (N f ) than the as-cast Ti-64-xB alloys. [3] The cast-and-extruded Ti-64-xB alloys also significantly outperformed their as-cast counterparts in terms of the creep resistance. [10] However, the effect of PM processing, which affects potential cost and performance advantages, on the microstructure and mechanical behavior has yet to be fully explored. The current work focused on the microstructure and elevated-temperature [728 K (455 C)] tensile and fatigue deformation behavior of Ti-64-1B through four different processing routes: PM-rolled, PM-extruded, IM-cast, and cast-and-extruded. In addition, the tensilestrength anisotropy and fatigue-creep interaction were explored VOLUME 42A, OCTOBER 2011 METALLURGICAL AND MATERIALS TRANSACTIONS A

2 II. EXPERIMENTAL A. Materials Processing The prealloyed Ti-64-1B powder was made by induction skull melting followed by inert gas atomization at Crucible Research Corporation (Pittsburgh, PA). The composition of the prealloyed powders as well as the final bulk products was measured through inductively coupled plasma optical emission spectroscopy and inert gas fluorescence. The Ti-64-1B powder was classified as 35 mesh (<500 lm in diameter). [22] In fabricating the PM-rolled Ti-64-1B plate, the prealloyed Ti-64-1B powder was degassed and sealed at RT for 24 hours and then at 573 K (300 C) for 24 hours. Hot isostatic pressing was performed at 1291 K (1018 C), 103 MPa for 3 hours on the powders, followed by heat treatment at 1573 K (1300 C) for 6 hours. The rolling process was uniaxial (three rounds, eight passes for each round, and before each round, the alloy billet was heated to 1373 K (1100 C). The dimension of the final rolled sheet was 96.5-cm long, 76-cm wide, and 2.5-cm thick. For the PM-extruded material, the same prealloyed Ti-64-1B powder was used. One kilogram of powder was packed inside a thick-walled Ti-64 can of 70-mm diameter and 130-mm length, vacuum outgassed at 573 K (300 C) for 24 hours, and sealed. The can was heated to 1473 K (1200 C), soaked for 1 hour, and compacted in an extrusion chamber using a blind-die. The compact was held at a pressure of 1400 MPa for 180 seconds followed by air cooling to RT. The compact was then extruded at 1373 K (1100 C) through a conical die using a 16:1 extrusion ratio. The final diameter of the extruded rods was approximately 2.0 cm. Castings of 7-cm diameter and 50-cm length were produced at Flowserve Corporation (Dayton, OH) via induction skull melting. The casting was hot isostatically pressed at 1173 K (900 C) and 100 MPa for 2 hours to remove porosity. This condition represented the as-cast state. A portion of the casting was sectioned and canned in commercially pure Ti in preparation for extruding. The extrusions were performed in the b-phase field at 1373 K (1100 C) using a 12:1 reduction ratio, and the strain rate was 1 s 1. The final diameter of the extruded rods was approximately 2.0 cm. No postprocessing heat treatments were performed for any of the studied materials. scanning electron microscope (FEI, Inc., Hillsboro, OR), or a JEOL* 6500 field emission scanning electron *JEOL is a trademark of Japan Electron Optics Ltd., Tokyo. microscope (JEOL Inc.). The average equiaxed grain size and a- and b-lath widths of the studied alloys were determined using the mean line intercept method. The lath widths of each alloy were measured perpendicular to the longitudinal axis of the laths for over 300 laths for each of the a and b phases. The average diameters (over 300 grains were measured) were multiplied by 1.74 to obtain the average equiaxed a-phase grain size. Electron backscatter diffraction (EBSD), performed using hardware and software manufactured by EDAX-TSL, Inc. (Mahwah, NJ), was used to evaluate the microstructures. Spatially-resolved EBSD maps were acquired at 20 kev using a step size ranging between 0.2 and 1.0 lm on samples polished through colloidal silica (0.06-lm average particle diameter). Transmission electron microscopy (TEM) analysis was performed using a Hitachi (Tokyo, Japan) HF-3300 transmission electron microscope/scanning transmission electron microscope at 300 kv. The samples were ground to a thickness of approximately 150 lm and then punched into 3-mmdiameter discs. The discs were then hand polished to approximately 100-lm thickness using 2400 grit silicon carbide paper and electropolished to perforation at a temperature of 15 C using a Struers (Westlake, OH) TenuPol-5 electropolishing unit at 28 V. The electropolishing solution contained 60 pct methanol, 33.5 pct 2-butoxyethanol, 6 pct perchloric acid, and 0.5 pct glycerin. In order to determine the weight percentages of Ti, Al, V, and B in the TiB, a phase, and b phase, microprobe analysis was conducted on mounted and polished samples using a JEOL JXA-8200 WD/ED combined microanalyzer operated at 10 kv and an emission current of approximately 55 na. The low accelerating voltage used was chosen to avoid overlap from the a, b, and TiB phases. The standards used were TiB 2 for Ti and B and pure Al and V for Al and V respectively. The crystals used were LiFH for V and Ti, TAP for Al, and a LDE 2 for B. X-ray diffraction (XRD) 2h scans were collected from 20 to 100 deg at a step size of deg using an X Pert Pro MPD diffractometer (PANalytical Inc., Almelo, Netherlands) with copper K a radiation. B. Microstructural Characterization The as-processed materials were prepared for imaging using conventional metallography techniques, and scanning electron microscopy (SEM) was used to examine the grain size, the a- and b-lath widths, the phase distributions, and their morphologies. Phase volume percents were determined using ImageJ image analysis software on several backscatter electron (BSE) SEM photomicrographs acquired using a CamScan44FE field emission scanning electron microscope (Tescan Inc., Brno, Czech Republic), an XL-30 field emission C. Tensile Experiments Tensile tests were performed on each of the alloys at a temperature of 728 K (455 C) at a strain rate of 10 3 s 1 using a servohydraulic testing machine described elsewhere. [23] This machine was also used to perform the fatigue tests, and the sample preparation was identical. Prior to testing, the 125-mm-long samples were electrodischarge machined (EDM) into a flat, dogbone geometry, and the EDM recast layers were removed by silicon carbide paper grinding through a 600 grit finish. Strain was measured during the tensile tests METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 42A, OCTOBER

3 with an alumina-arm extensometer attached directly to the heated gage section of the sample. The samples were allowed to soak at the desired temperature for 15 minutes prior to tensile testing. In order to evaluate the tensile-strength anisotropy of the extruded and rolled materials, dogbone samples were EDM cut along the transverse orientation. Ni-based superalloy (IN 718) grips were EDM machined, and an Ernest Fullam tensile stage (MTI Instruments, Albany, NY) was used to perform these experiments. A description of the tensile testing technique involving use of this stage is provided elsewhere. [17] A minimum of two tensile tests were performed for each tensile condition examined, and the average properties were reported. D. Fatigue Experiments Open-air, tension-tension (R = 0.1) fatigue experiments were performed at a temperature of 728 K (455 C) using a test frequency of 5 Hz. The maximum applied stresses for the fatigue tests varied between 300 and 700 MPa. Strain was monitored using an aluminarod high-temperature extensometer, with a 12-mm gage length, attached to the gage section of the specimen. The strain-life behavior was documented throughout the experiments, and both the E values and hysteresis behavior were tracked at certain cycles during the experiments. In such cases, 20 stress-strain data points were acquired throughout the desired cycle. The specimens were locally heated using a Barber Coleman temperature controller and two sets of heating banks, each containing four evenly spaced quartz lamps, which were located approximately 5 mm above and below the sample. Specimen temperatures were monitored by four chromel-alumel type-k thermocouples located within the specimen s gage section. Targeted test temperatures were maintained within ±5 C. The test specimens were soaked at 728 K (455 C) for at least 20 minutes prior to applying load in order to minimize the thermal stresses. At least two repeat experiments were performed for each maximum stress level examined. For the fractured samples, the furnace was turned off immediately upon fracture to prevent oxidation of the exposed surfaces. The fractured surfaces and gage sections of the deformed samples, for both samples taken to failure and runout samples, were evaluated using SEM. Several of the specimens, which exhibited runout (N>1,000,000 cycles without failing), were tensile tested at 728 K (455 C) in order to evaluate if low-stress fatigue degraded the material. III. RESULTS A. Microstructure Table I lists the compositions of the studied alloys. Table II lists the average volume percents of the TiB and b phases along with the equiaxed a grain size and average a-andb-lath widths. Photomicrographs of the as-processed microstructures are illustrated in Figures 1 through 3. The TiB phase is dark, the a phase is gray, and the b phase is light. The crystal structures of the a- Ti, b-ti, and TiB phases are known to be hexagonalclosed-packed (hcp), body-centered-cubic (bcc), and orthorhombic, respectively. The volume percent of TiB phase in each alloy was approximately 6. The volume percent of the b phase was approximately 15 pct. The average a-lath width in the PM-rolled condition was approximately 2.7 lm, which was between that of the as-cast (3.6 lm) and cast-and-extruded (1.0 lm) conditions. The average grain size of the equiaxed a phase was 3.6 lm in the PM-extruded condition, which was significantly smaller than the a-grain size (7.4 lm) in the PM-rolled condition. The microprobe data (Table III) showed that the b phase was enriched with V and the a phase was enriched with Al and Ti. There was no detectable amount of B in either the a or b phase, and the TiB phase exhibited slightly more B than Ti when Table I. Compositions of the Studied Ti-64-1B Alloys Alloy Processing Al (Wt Pct) V (Wt Pct) B (Wt Pct) Fe (Wt Pct) O (Wt Pct) N (Wt Pct) C (Wt Pct) H (Wt Pct) Ti-64-1B pre-alloyed powder Ti-64-1B as-cast Ti-64-1B cast-and-extruded N/A *Ti was the balance in the materials. Table II. TiB and b-phase Volume Percent (V p ), a- and b-lath Widths (k), and a Grain Size (d a ) Alloy Processing V p TiB* V p b* k a (lm) k b (lm) d a (lm) Ti-64-1B PM-rolled 5.9 ± ± ± ± ± 0.7 Ti-64-1B PM-extruded 6.3 ± ± 1.8 NA NA 3.6 ± 0.8 Ti-64-1B as-cast 5.8 ± ± ± ± 0.2 NA Ti-64-1B cast-and-extruded 6.0 ± ± ± ± 0.1 NA *The a phase was the balance in the microstructure. NA: not available VOLUME 42A, OCTOBER 2011 METALLURGICAL AND MATERIALS TRANSACTIONS A

4 Fig. 1 BSE SEM images of the PM-rolled Ti-64-1B alloy microstructures in the (a) longitudinal and (b) transverse orientations. compared on an atomic percent basis. There was a small amount of V (approximately 3.1 to 3.5 wt pct) detected in the TiB phase. EBSD analysis indicated that the a and TiB phases for the PM-extruded material were textured (Figure 4). The TiB phase was preferentially oriented with [020] parallel to the extrusion axis, while the basal plane of the a phase was oriented perpendicular to the extrusion axis, preferentially. The texture of the a phase in the cast-andextruded condition was similar to that of the PMextruded material (Figure 5). For the PM-rolled material, the a phase was not strongly textured (Figure 6). Figure 6(c) indicates that the TiB phase was also not strongly textured in the PM-rolled condition. Figure 7 illustrates XRD intensity vs 2h plots for the studied materials. The plots for the PM-rolled sheet (Figures 7(a) through (c)) were acquired for the three orientations of the rolled plate, namely, face (F), longitudinal (L), and transverse (T). The analysis suggested that a strong texture was not developed during the rolling process. The as-cast material was not textured (Figure 7(d)). The greatest peak for the PMextruded and cast-and-extruded conditions was for the basal plane (0002). Thus, the extrusion aligned the basal plane perpendicular to the extrusion axis for these two extruded materials. The largest peak for the retained b phase was for the (110) plane. This (110) b-phase peak Fig. 2 BSE SEM images of the PM-extruded Ti-64-1B alloy microstructures in the (a) longitudinal and (b) transverse orientations. In (a), the TiB phase was aligned in the extrusion direction (horizontal), and in (b), the TiB phase is pointing out of the plane of the page. has also been found in cast-and-extruded Ti-64 and Ti B alloys. [3] It is believed that the b-phase field extrusion procedure caused (110) to be oriented perpendicular to the extrusion axis, similar to that developed in a-fe (bcc) by the cold-drawn process. [24,25] Due to the Burgers orientation relationship between a and b Ti, [26] the basal plane of a phase is expected to orient itself parallel to the (110)b phase during the cooldown from the b-phase field. B. Tensile Behavior The average 728 K (455 C) tensile properties are listed in Table IV, and Figure 8 illustrates the 728 K (455 C) tensile stress vs strain behavior of representative samples. The YS and UTS values of the PM-rolled and as-cast materials were significantly lower than the PM-extruded and cast-and-extruded materials. The YS of the PM-rolled material was similar to that of the ascast material, while its e f value was much higher. The YS, UTS, and e f values of the PM-extruded material were similar to those of the cast-and-extruded material. The e f value for the as-cast material was relatively low (2.5 pct), while each of the other alloys exhibited a minimum e f value of 8.2 pct. METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 42A, OCTOBER

5 Tensile tests in the T orientation of the extrusions and plates were performed to assess the effect of texture on the tensile properties. The average 728 K (455 C) YS and UTS values in both the L and T orientations are shown in Figure 9(a), while the RT properties are shown in Figure 9(b). The 728 K (455 C) YS and UTS values were roughly 60 pct of those at RT. For the PM-rolled Fig. 3 BSE SEM images of the IM processed Ti-64-1B alloy microstructures: (a) as-cast and (b) cast-and-extruded. In (b), the TiB phase was aligned in the extrusion direction (horizontal). condition, the tensile properties in the L and T orientations were similar. This was expected because this material was not strongly textured. However, for the extruded materials, the YS and UTS values in the L orientation were greater than those in the T orientation. This anisotropy was expected due to the texture caused by the extrusion procedure. The tensile anisotropy was much stronger for the PM-extruded material than for the cast-and-extruded material. The UTS difference between L and T orientations for the PM-extruded material was 97 MPa [728 K (455 C)] and 167 MPa (RT) (Figure 9); the YS difference was 58 MPa [728 K (455 C)] and 82 MPa (RT). For the cast-and-extruded material, the UTS difference was 35 MPa [728 K (455 C)] and 73 MPa (RT); the YS difference was 35 MPa [728 K (455 C)] and 52 MPa (RT). In addition, tensile tests in the T orientation of the cast-and-extruded Ti-64 and Ti B alloys were also performed to evaluate the effect of B on the tensile anisotropy. Similar to the PM-extruded and cast-andextruded Ti-64-1B alloys, these two alloys were also strongly textured with the basal plane of a phase perpendicular to the extrusion axis. [3] The YS and UTS in both the L and T orientations are shown in Figures 10(a) [728 K (455 C)] and 10(b) (RT). It is clear that the tensile strength in the L orientation for each of the Ti-64-xB cast-and-extruded alloys was greater than that in the T orientation. It is evident that the tensile-strength anisotropy was more pronounced for the cast-and-extruded Ti-64 and Ti B alloys compared to the Ti-64-1B alloy. Thus, the tensilestrength anisotropy was weakened by B addition. It is noteworthy that the UTS and YS values of the Ti-64-1B and Ti B cast-and-extruded alloys and the Ti-64-1B PM-extruded alloy in the T orientation were always larger than those for the as-cast Ti-64-xB alloys. [3] Thus, although the extruded Ti-64-xB alloys exhibited anisotropic behavior, the B additions in the extruded alloys significantly improved the tensile properties independent of sample orientation. During the in-situ tensile tests for the extruded Ti-64-xB alloys in the T orientation, debonding was not observed, indicating a strong interface bonding between the TiB phase and the Ti-64 Table III. Phase Compositions in Weight Percent for the Studied Ti-64-1B Alloys TiB 77.4 ± 0.2 (46.5 ± 0.3) 0.1 ± 0.1 (0.1 ± 0.1) 3.1 ± 0.1 (1.8 ± 0.1) 19.4 ± 0.2 (51.7 ± 0.2) a 89.5 ± 0.2 (84.6 ± 0.2) 7.6 ± 0.2 (12.8 ± 0.2) 2.8 ± 0.3 (2.6 ± 0.3) 0 (0) b 81.8 (79.6) 4.7 (8.1) 13.5 (12.3) 0 (0) Ti-64-1B as-cast TiB 77.8 ± 0.5 (47.4 ± 0.6) 0.06 ± 0.05 (0.1 ± 0.1) 3.4 ± 0.1 (1.9 ± 0.04) 18.7 ± 0.4 (50.4 ± 0.7) a 89.2 ± 0.2 (84.4 ± 0.6) 8.0 ± 0.1 (12.8 ± 1.1) 2.9 ± 0.06 (2.8 ± 0.5) 0 (0) b 76.9 ± 0.7 (75.9 ± 0.5) 3.2 ± 0.2 (5.6 ± 0.4) 19.9 ± 0.9 (18.5 ± 0.9) 0 (0) Ti-64-1B Alloy Processing Phase Ti* Al* V* B* Ti-64-1B PM-rolled TiB 78.1 ± 0.7 (47.7 ± 1.0) 0.2 ± 0.2 (0.2 ± 0.2) 3.2 ± 0.1 (1.8 ± 0.1) 18.5 ± 0.6 (50.2 ± 1.0) a 89.4 ± 0.1 (84.4 ± 0.1) 7.9 ± 0.04 (13.2 ± 0.1) 2.7 ± 0.1 (2.4 ± 0.1) 0 (0) b 78.2 ± 0.4 (76.9 ± 0.3) 3.5 ± 0.1 (6.2 ± 0.2) 18.3 ± 0.5 (16.9 ± 0.5) 0 (0) Ti-64-1B PMextruded cast-andextruded *Atomic percent is in parentheses. TiB 78.3 ± 0.5 (48.5 ± 1.5) 0.1 ± 0.1 (0.1 ± 0.1) 3.5 ± 0.2 (2.0 ± 0.2) 17.9 ± 1.2 (49.1 ± 2.2) a 88.7 ± 0.4 (84.2 ± 0.3) 7.2 ± 0.3 (12.2 ± 0.4) 4.2 ± 0.7 (3.7 ± 0.6) 0 (0) b 84.2 (81.2) 5.5 (9.4) 10.4 (9.4) 0 (0) 3050 VOLUME 42A, OCTOBER 2011 METALLURGICAL AND MATERIALS TRANSACTIONS A

6 Fig. 5 EBSD data from a representative scan of the cross section (the orientation out of the page was the extrusion direction) for a cast-and-extruded Ti-64-1B sample: (a) a-phase orientation map; (b) a-phase (0001) pole figure; and (c) (001), (010), and (001) pole figures for the orthorhombic B27 TiB phase. Fig. 4 EBSD data from a representative scan of the cross section (the orientation out of the page was the extrusion direction) for a PM-extruded Ti-64-1B sample: (a) a-phase orientation map; (b) a- phase (0001) pole figure; and (c) (001), (010), and (001) pole figures for the orthorhombic B27 TiB phase. matrix. This phenomenon is consistent with the observations of Gorsse and Miracle. [6] C. Fatigue Behavior The fatigue S-N curves obtained for the studied Ti-64-1B alloys are shown in Figure 11. The PM-extruded material exhibited significantly larger N f values than all other studied materials. The PM-rolled material exhibited similar fatigue lives to those of the cast-and-extruded material. The N f values of the as-cast Ti-64-1B were the lowest. TiB-phase cracking was evident within the deformed gage sections (Figure 12). This behavior was similar to that for a Ti-64/TiB alloy containing 0.5 wt pct B. [14] For the fatigue samples that fractured, stress vs strain plots during defined fatigue cycles are shown in Figure 13. None of the PM samples exhibited an E-value reduction during the experiments (Figures 13(a) and (b)), which suggested that an insignificant amount of fatigue cracking or fatigue crack growth occurred. Selected fatigue runout samples (N > 1,000,000) were tensile tested at 728 K (455 C), as shown in Table V. Comparing these results to those in Table IV, it is evident that there was no tensile strength loss at 728 K (455 C) due to fatigue for the low-stress fatigue samples METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 42A, OCTOBER

7 cases, the measured volume percents of the TiB, a, and b phases were similar both before and after the experiments. The stress vs strain plots during defined fatigue cycles for the as-cast and cast-and-extruded Ti-64-xB alloys were described previously. [3] An E-value decrease was not observed in the runout samples for these alloys either. However, for the samples that fractured, a slight E-value decrease was observed, which was consistent with the TiB-phase cracking occurring during these experiments. The fracture initiation sites of each material tested were typically located at surface locations. In some cases, multiple surface crack initiation sites were evident. Fig. 6 EBSD data from a representative scan of the cross section (the orientation out of the page was the rolling direction) for a Ti- 64-1B PM-rolled sample: (a) a-phase orientation map; (b) a-phase (0001) pole figure; and (c) (001), (010), and (001) pole figures for the orthorhombic B27 TiB phase. that experienced runout. The increasing strain with increasing cycle number suggests that ratcheting occurred during the fatigue tests. No phase instability was observed on the postfatigue gage sections for any of the samples tested even after more than 2.1 million cycles of fatigue exposure at 728 K (455 C). In such IV. DISCUSSION A. Microstructure The PM-rolled material exhibited a duplex microstructure, which contained both lenticular and equiaxed a phase. During the rolling process, the billet surface temperature dropped to 1223 K (950 C), which is below the b-transus temperature (999 ± 14 C). [27] The rolling process following by the 1033 K (760 C) annealing treatment likely introduced some equiaxed a phase into the microstructure. The different thermomechanical processing routes caused the a-phase morphology difference in the PMrolled and PM-extruded materials. In the PM-extruded material, the equiaxed a phase was predominant with the b phase decorating the grain boundaries (Figure 2). The a grain diameter (3.6 lm) was only about half that for the PM-rolled material (7.4 lm). The processing of the PM-extruded material was identical to that for a PM-extruded Ti B alloy, [28] where the Ti B prealloyed powder exhibited a martensite microstructure. After blind-die compaction at 1473 K (1200 C), the alloy exhibited equiaxed a phase, [28] and after extrusion at 1373 K (1100 C), the alloy revealed a fine-grained (~5 lm) equiaxed a phase, which is similar to that of the PM-extruded Ti-64-1B alloy in the current study. The a-phase morphology of the current PMextruded alloy is also similar to that for the PM Ti B alloy studied by Bhat et al. [29] Since the as-cast and cast-and-extruded materials were both cooled from above the b-transus temperature and no thermomechanical processing was performed thereafter, the resulting a-phase morphology was fully lenticular. A finer prior-b grain size and a finer a-colony size were exhibited by the extruded material. The average prior-b grain size was approximately 200 lm in the as-cast material and 30 lm in the cast-andextruded material. The average a-colony size was approximately 30 lm in the as-cast material and 10 lm in the cast-and-extruded material. The extrusion procedure also resulted in a significantly refined lath width. In fact, the a-lath width for the cast-and-extruded material was 3.6 times finer than that for the as-cast material (Table II). This was likely a result of the extrusion procedure followed by faster cooling than that for the as-cast melt. The dislocations produced during the extrusion process may act as effective sites for 3052 VOLUME 42A, OCTOBER 2011 METALLURGICAL AND MATERIALS TRANSACTIONS A

8 Fig. 7 XRD intensity vs 2h plots for the studied Ti-64-1B alloys: (a) through (c) PM-rolled F, L, and T orientations; (d) as-cast; (e) PM-extruded T orientation; and (f) cast-and-extruded T orientation. Note the ratio of (0002) peak and ð1011þpeak in (e) and (f). This indicated that the extruded materials were strongly textured with (0002) perpendicular to the extrusion axis, while the PM-rolled material was not strongly textured. Table IV. Average 728 K (455 C) Tensile Properties for the Studied Ti-64-1B Alloys Alloy Processing E (GPa) YS (MPa) UTS (MPa) e f (Pct) Ti-64-1B PM-rolled Ti-64-1B PM-extruded Ti-64-1B as-cast Ti-64-1B cast-and-extruded heterogeneous nucleation of the a laths during cooling. The enrichment of V in the b phase and Al in the a phase (microprobe analysis in Table III) was expected, since V and Al are strong b-phase and a-phase stabilizers, respectively. [30] For the PM-extruded material, the extrusion process aligned the TiB phase in the extrusion direction with [020] aligned parallel to the extrusion direction (Figure 2). Thermomechanical processing in the b-phase field allows reasonably easy plastic flow and facilitates rigid body rotation of the intermetallic TiB, causing alignment along the direction of flow. [31] The alignment of the TiB phase was also observed for a Ti B alloy using the same PM processing conditions. [28] Fig. 8 The 728 K (455 C) tensile stress vs strain behavior of representative Ti-64-1B alloys for each processing condition. For the PM-extruded material, the basal plane of the a phase was predominantly oriented perpendicular to the extrusion axis. This texture was similar to that METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 42A, OCTOBER

9 Fig. 9 Comparison of YS and UTS in L orientation (plotted with columns) and T orientation (plotted with cylinders) for the PM-rolled, PM-extruded, and cast-and-extruded Ti-64-1B alloys: (a) 728 K (455 C) properties and (b) RT [295 K (22 C)] properties. *Data taken from Ref. 12. Fig. 10 Comparison of YS and UTS in L orientation (plotted with rectangular columns) and T orientation (plotted with cylinders) for the cast-and-extruded Ti-64-xB alloys; (a) 728 K (455 C) properties and (b) RT [295 K (22 C)] properties. resulting from hot rolling of Ti-64 rods. [32,33] During the b-phase field extrusion, it is expected that the b phase will preferentially align with the {110} planes parallel to the extrusion direction as the b phase has 12 slip systems associated with the {110} family of planes. It was Fig. 11 Maximum applied stress vs cycles-to-failure curves for the studied Ti-64-1B alloys. Runout specimens are marked with arrows. previously noted that a strong (110) texture was found in cold-drawn a-fe (bcc). [24,25] The {110} planes are located either 0, 60, or 90 deg apart in the cubic system. The Burgers orientation relationship between a and b is known to be (110)b//(0002)a; ½111Šb//½1120Ša. [26] Following transformation to the a phase, the texture of the b phase would be expected to be inherited by the a phase in a manner consistent with the Burgers relationship, as has been observed previously for Ti-64. [34] Thus, after the transformation to the a phase, which occurred during the cooldown from the b-phase field, the basal plane normal of the a phase is expected to be either 0, 60, or 90 deg from the extrusion direction, which is in agreement with the experimental observations. In addition, the largest peak for the retained b phase was for the (110) plane. This was also observed in previous studies of Ti-64. [32,33] Based on the orientation relationship and the symmetries of the cubic and hexagonal structures, a total of 12 crystallographically distinct variants of the a phase can result from the transformation of a single b grain. Variant selection would dictate how much basal plane a is 0, 60, or 90 deg oriented. The hcp-to-bcc phase transformation in Ti was observed in situ by EBSD [35] and by neutron diffraction techniques. [36] Lonardelli et al. [36] studied the texture evolution of b fi a phase transformation in Ti-64 alloy by insitu neutron diffraction and suggested that the growth of a domains is controlled by high-temperature a orientations inherited from the b grains with the Burgers orientation relation. The mechanism of the a-phase texture development in the cast-and-extruded material is believed to be the same as that in the PM-extruded material. The peak ratio between the (0002) and the ð1011þplanes was greater in the PM-extruded material (3.4) than in the cast-and-extruded material (1.1), as shown in Figures 7(e) and (f). This suggests that the (0002) texture of the a phase in the PM-extruded material was stronger than that in the cast-and-extruded material VOLUME 42A, OCTOBER 2011 METALLURGICAL AND MATERIALS TRANSACTIONS A

10 Fig. 12 BSE SEM photomicrographs taken from subsurface sections close to the fatigue fracture surface of the studied Ti-64-1B alloys: (a) PM-rolled sample (N f = 266,638 cycles, r max = 550 MPa); (b) PM-extruded sample (N f = 162,628 cycles, r max = 650 MPa); (c) as-cast sample (N f = 32,718 cycles, r max = 350 MPa); and (d) cast-and-extruded sample (N f = 79,230 cycles, r max = 500 MPa). TiB cracking is evident. B. Tensile Behavior Starting with the same prealloyed Ti-64-1B powder, the extrusion process resulted in a significantly stronger material than the rolling process. As the vast majority of the microstructure consisted of the a phase, the texture of the a phase is an important microstructural parameter in terms of the mechanical properties. If the basal plane is either normal or parallel to the extrusion/tensile axis, then slip along the basal plane is prevented. Thus, this represents a hard orientation. Gorsse and Miracle [6] studied the influence of TiB phase orientation on the tensile properties of PM-processed Ti-64/TiB at RT and 573 K (300 C). Their results show that the tensile strength and E values of the composites with aligned TiB phase were always higher than composites with randomly-oriented TiB phase. Thus, in the current study, both the alignment of the TiB phase and the a-phase texture in the PM-extruded Ti-64-1B alloy contributed to its higher tensile strength and E values compared with those for the PM-rolled material. In addition, the smaller a grain size in the PM-extruded material also contributed to its higher strength. [37] The E value of the PM-extruded material (117 GPa at 728 K (455 C) and 142 GPa at RT) was higher than that of the cast-and-extruded material (113 GPa at 728 K (455 C) and 128 GPa at RT). This was expected to be a result of the stronger (0002) a-phase texture in the PM-extruded material compared with the cast-andextruded material. Although the a-phase morphology in the PM-rolled (duplex) and as-cast (lenticular) materials was different, the tensile strength of each material was similar (Figure 8). This result is not unexpected since strength is virtually unaffected by the morphology of the a phase. [38] In addition, the tensile properties of the PMextruded (equiaxed) material were similar to those of the cast-and-extruded (lenticular) material (Figure 8), although the a-phase morphologies in these two alloys were significantly different. In addition, the E values of these two alloys were quite close (Table IV). In general, duplex microstructures exhibit higher e f values compared with lenticular microstructure. [37] This may be one reason why the PM-rolled Ti-64-1B exhibited a higher e f than the as-cast Ti-64-1B. The tensile-strength anisotropy (Figures 9 and 10) in the extruded materials resulted from the texture caused by the extrusion procedure. At 728 K (455 C), the UTS difference between L and T orientations decreased from 103 MPa fi 87 MPa fi 35 MPa as the B content increased from 0 wt pct fi 0.1 wt pct fi 1 wt pct METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 42A, OCTOBER

11 Table VI. Tensile Properties for the As-Cast Ti-64-xB Alloys Alloy Temperature E (GPa) YS (MPa) UTS (MPa) e f (pct) Ti K (455 C) Ti-64 [2] RT Ti-64-1B 728 K (455 C) Ti-64-1B [2] RT Fig. 13 Stress vs strain plots at the marked cycle numbers for the studied Ti-64-1B alloys at 728 K (455 C): (a) PM-rolled and (b) PM-extruded. The slope of each curve is indicated by the listed Young s modulus value. Alloy Table V. Average 728 K (455 C) Tensile Properties for the Runout (N > 1,000,000) Ti-64-1B Alloys Processing E (GPa) YS (MPa) UTS (MPa) e f (Pct) Ti-64-1B PM-rolled Ti-64-1B PM-extruded Ti-64-1B as-cast Ti-64-1B cast-and-extruded (Figure 10). The UTS difference at RT followed the same trend. Similar to the UTS trend, the YS at 728 K (455 C) as well as at RT also decreased with the increase of B content (Figure 10). This trend is reasonable since the XRD data (Figure 14 and Figure 7(f)) indicated that the relative intensity of the (0002) peak Fig. 14 XRD intensity vs 2h plots for transverse sections of castand-extruded (a) Ti-64 and (b) Ti B. was weakened by increased B content. This was also reinforced by XRD pole figures data for the cast-andextruded Ti-64 and Ti-64-1B alloys. [3] For the cast-and VOLUME 42A, OCTOBER 2011 METALLURGICAL AND MATERIALS TRANSACTIONS A

12 extruded Ti-64-1B alloy, the TiB phase tended to compensate for the loss of strength in the T orientation due to the weakening texture of a phase. Thus, the combined effect of both the TiB-strengthening and the a-phase texture weakening resulted in the lower tensilestrength anisotropy in the cast-and-extruded Ti-64-1B alloy compared to the cast-and-extruded Ti-64 and Ti B alloys. For the PM-extruded material, the tensile-strength anisotropy was more pronounced. The reason for this result is that the a-phase texture in the PM-extruded material was more pronounced: the ð0002þ=ð1011þpeak height ratio was 3.4 (Figure 7(e)) compared with 1.1 (Figure 7(f)) for the cast-and-extruded material. In summary, the tensile strength in the T orientation for the extruded materials is lower than in the L orientation. However, it is still higher than the tensile strength of the as-cast Ti-64-1B and Ti-64 alloys at both 728 K (455 C) and RT (Table VI). The B addition and processing did not degrade the tensile properties in either the L or the T orientation compared with the baseline as-cast Ti-64 alloy. Thus, the extruded Ti-64-xB alloys have the potential to replace Ti-64 for both elevated-temperature and RT structural applications driven by tensile and fatigue conditions. C. Fatigue Behavior Ti-64 exhibits superior high-cycle smooth bar fatigue life when the slip length is small. [39 41] Since a grain boundary can act as an effective obstacle to dislocation motion, reducing the as-cast grain size of Ti-64 alloy through B addition will reduce the dislocation slip length. This effect was confirmed in the as-cast Ti-64-1B alloy. Figure 15 illustrates a larger dislocation density at the grain boundary area of a fatigue-tested as-cast Ti-64-1B specimen. Figure 15(c) illustrates the dislocation pileup at the grain boundary. Relatively, large dislocation densities were observed in areas close to grain boundaries. As was shown in Figure 11, the PM materials exhibited greater fatigue resistance than the IM materials. The PM-extruded material exhibited the greatest fatigue life. There are several microstructural features that contributed to the improved fatigue resistance exhibited by the PM materials. It is well established that low-cycle fatigue (LCF) behavior is favored by an equiaxed a morphology, [38] and smaller grain size (equiaxed a) results in better fatigue properties. [37] Reducing the equiaxed a grain size in Ti-64 alloy from 12 to 2 lm increased the fatigue strength from 560 to 720 MPa. [42,43] The equiaxed a grain size in the PMextruded material (3.6 lm) was significantly finer than that in the PM-rolled material (7.4 lm). Thus, PM processing induced a finer equiaxed a grain and this contributed to its enhanced fatigue lives compared to those of the IM materials. The a-phase texture (basal plane perpendicular to tensile axis) enhanced the fatigue resistance in the PM-extruded material. The texture of the anisotropic a phase is critical to its mechanical behavior. At elevated temperatures, the a phase deforms due to a and some c + a slip, where a slip is always Fig. 15 TEM bright-field image of a fatigue-tested as-cast Ti-64-1B alloy (N f = 14,120 cycles, r max = 450 MPa). (a) Grain boundary area (TiB phase can be seen); (b) higher-magnification image in the solid square of (a), where dislocations are evident; and (c) highermagnification image in the broken square of (a), dislocations piled up at the grain boundary (G1: grain 1; G2: grain 2). METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 42A, OCTOBER

13 Fig. 16 Temperature dependence of CRSS for slip with a and c + a vectors in single crystals of Ti-6.6Al [44]. dominant for a-ti. [44] Paton et al. [45] investigated the temperature dependence of the critical resolved shear stress (CRSS) for slip with a and c + a Burgers vectors in single crystals of Ti-6.6Al (Figure 16). It was shown that at 728 K (455 C), the CRSS for c + a slip is approximately 400 MPa, which is significantly higher than that for a slip (100 MPa). [45] When the basal plane is perpendicular to the tensile axis, as was the case for the extruded materials, a slip is prevented. Thus, deformation is only possible in c + a slip for this case. As has been described earlier, the (0002) texture of the a phase in the PM-extruded material was more pronounced than that in the cast-and-extruded material. This more pronounced texture led to less dislocation activity and likely enhanced the fatigue resistance of the PM-extruded material. Figure 17 illustrates an area where dislocations were constrained between the two TiB whiskers in a fatigue-tested PM-rolled specimen. The a-phase texture of the cast-and-extruded material was expected to be the primary reason for its increased fatigue strength compared with the as-cast material. The average a-colony size in the cast-and-extruded material is estimated to be 10 lm, compared with 30 lm for the as-cast material. In addition, the a-lath width for the cast-and-extruded material was 3.6 times finer than that of the as-cast material (Table II). Thus, the reduction of a-colony size and lath width is expected to have also contributed to improved fatigue resistance. [37] D. Ratcheting Behavior Strain ratcheting or cyclic creep describes the phenomenon of progressive damage accumulation under cyclic loading. [46] The strain vs time plots for the PMrolled material are provided in Figure 18. The total strain was taken at the peak stress of certain cycle numbers ( ,1910 5,2910 5,3910 5, ) and plotted with the elapsed time. This curve closely resembles a typical creep curve exhibiting primary and secondary creep stages. Similar to those in conventional creep tests, higher stress levels resulted in higher strain Fig. 17 TEM bright-field image of a fatigue-tested PM-rolled Ti- 64-1B alloy (N f = 12,998 cycles, r max = 600 MPa). Dislocations were evident between two TiB whiskers. rates. Comparing the total strain vs time curves from the fatigue tests with the creep strain vs time curves from the conventional creep tests at the same temperature and same maximum stress [10] (Figure 18), it was observed that the total strain during fatigue was always lower than the creep strain. Although the fatigue sample represented in Figure 18 only experienced the peak stress once in a cycle, the minimum strain rate during fatigue ( /s) was close to that for the conventional creep test ( /s). [47] The total strain vs time curves at a maximum stress level of 500 MPa are compared for the PM-processed materials in Figure 19. It is evident that the more fatigue resistant material (PM extruded) exhibited less strain ratcheting. The cast-and-extruded Ti-64-xB alloys exhibited similar behaviors to those of the PM-processed materials (Figure 20). In this case, the strain rates taken from the fatigue experiments were slightly lower than the creep strain rates. For the cast-and-extruded Ti-64 alloy at 450 MPa, the minimum total strain rate during fatigue was s 1, compared with a minimum creep strain rate of s 1 ; [47] for the cast-andextruded Ti B alloy at 450 MPa, the minimum strain rate during fatigue was s 1, compared with a minimum creep strain rate of s 1. [10] The total strain during the fatigue tests was always lower in the Ti-64 alloy than in the Ti B alloy (Figures 20(a) and (b)). This phenomenon was consistent with conventional creep tests, because the cast-and-extruded Ti-64 alloy was more creep resistant than the cast-and-extruded Ti B alloy. [10] For the cast-and-extruded Ti-64-1B alloy, the minimum total strain rate during fatigue was s 1, compared with a minimum creep strain rate of s 1. [10] Another typical example is the ascast Ti B alloy (Figure 21). (This sample broke at 3058 VOLUME 42A, OCTOBER 2011 METALLURGICAL AND MATERIALS TRANSACTIONS A

14 Fig. 19 Strain vs time plot during fatigue tests at 728 K (455 C), 500 to 50 MPa for the PM-processed Ti-64-1B alloys. Both samples exhibited runout (N f > 1,000,000). The minimum strain creep rates are labeled next to the corresponding curves. were subjected continuously to the same maximum stress levels. In the conventional creep tests, secondary creep starts when the dislocation substructure is stable. In the fatigue tests, the large amount of dislocations generated by fatigue could facilitate the creep deformation process, and it is possible that the dislocation substructure never becomes stable. In addition, defects induced by fatigue (such as microvoids) could also increase the strain rate. Fig. 18 Strain vs time plots for the PM-rolled Ti-64-1B alloy at 728 K (455 C): (a) fatigue test at 400 to 40 MPa (runout) and creep test at 400 MPa; [47] and (b) fatigue test at 450 to 45 MPa (runout) and creep test at 450 MPa [47]. 990,554 cycles, and the E value dropped from 88 GPa at N = 100 cycles to 76 GPa at N = 990,000 cycles.) For this sample, the total strain vs time curve eventually entered the tertiary regime before the final fracture of the sample, which was due to the severe fatigue damage that occurred in the material. The strain rates just before failure were similar in both creep ( s) and fatigue ( s 1 ). Thus, the fatigue specimen is believed to have failed due to the accelerated damage caused by strain ratcheting. In summary, the strain ratcheting during the loadcontrolled fatigue tests at 728 K (455 C) was significant. Even though the fatigue samples only experienced the highest stress once per cycle, the total strain rates were close to or even higher than the creep strain rates measured in conventional creep tests, where specimens V. SUMMARY AND CONCLUSIONS Ti-64-1B alloys, processed using PM and IM methods, were fatigue tested at 728 K (455 C) and maximum applied stresses between 300 and 700 MPa. The microstructure of PM-processed materials was significantly different from those for the as-cast and cast-and-extruded materials. The PM-rolled material exhibited a duplex microstructure and the PM-extruded material exhibited a fine, equiaxed microstructure. The a-phase morphology difference in the PM-rolled and PM-extruded materials was the result of different thermomechanical processing routes. The b-phase field extrusion aligned the {110} planes of b-phase perpendicular to the extrusion axis. After the b fi a phase transformation, a strong a-phase texture was developed in the PM-extruded and cast-andextruded materials. For the PM-extruded material, the finer equiaxed a grain size, the a-phase texture, and the alignment of TiB phase all contributed to its higher tensile strength and Young s modulus compared with the PM-rolled material. The a-phase texture in the extruded materials resulted in tensile-strength anisotropy. The YS and UTS for these alloys in the T orientation were lower than in the L orientation. As the a-phase texture became less pronounced with increasing B content in the cast-andextruded materials, the tensile-strength anisotropy METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 42A, OCTOBER

15 Fig. 21 Strain vs time plots for the as-cast Ti B alloy at 728 K (455 C). The N f value of the fatigue-tested sample was 990,554 cycles. Fig. 20 Strain vs time plots for the cast-and-extruded Ti-64-xB alloys at 728 K (455 C). weakened. However, the strength in the T orientation for the extruded Ti-64-1B and Ti B was always stronger than that for the Ti-64 as-cast alloy. At 728 K (455 C), the UTS values of cast-and-extruded Ti-64-xB alloys in the T orientation were at least 48 pct higher than the UTS of the baseline as-cast Ti-64 alloy (435 MPa). Thus, it was shown that the tensile strength of B-modified Ti-64 is always greater than that for Ti-64. The PM-processed materials exhibited greater fatigue resistance than the as-cast and cast-and-extruded materials. Comparing the current data with the data obtained in previous studies, [3,9] the PM-extruded Ti-64-1B alloy outperformed all the studied Ti-64-xB alloys in fatigue. The equiaxed a-phase introduced in the PM-rolled (duplex microstructure) and PM-extruded (equiaxed microstructure) alloys played a key role since the LCF resistance of Ti alloys is favored by an equiaxed a morphology. The fine, equiaxed grain size and the strong a-phase texture were both important factors for the greater fatigue resistance exhibited by the PMextruded material. For the IM processed materials, the cast-and-extruded Ti-64-1B alloy exhibited greater fatigue resistance than the as-cast Ti-64-1B alloy. This was explained to be a result of the significantly smaller grain size, a-colony size, and a-lath widths, as well as the pronounced a-phase texture exhibited in the cast-andextruded material. Strain ratcheting during the load-controlled fatigue tests at 728 K (455 C) was significant. Although the samples during the fatigue tests only experienced the maximum stress once in every cycle, in several cases, the total strain rate during the fatigue tests was close to the minimum creep strain rate from the conventional creep tests continuously subjected to the same maximum stress level VOLUME 42A, OCTOBER 2011 METALLURGICAL AND MATERIALS TRANSACTIONS A