The Effects of Forging and Rolling on Microstructure in O+BCC Tialnb

Materials Science and Engineering A279 (2000) 118–129 www.elsevier.com/locate/msea

The effects of forging and rolling on microstructure in O+BCC TiAlNb alloys

C.J. Boehlert *

Department of Mechanical Engineering, Johns Hopkins Uni6ersity, 3400 North Charles Street, Baltimore, MD 21218, USA

Received 1 July 1999; received in revised form 8 September 1999

Abstract

The effects of hot upset forging and hot pack rolling on microstructure of orthorhombic (O)+body-centered cubic (BCC) TiAlNb alloys was investigated. The starting materials were melted ingots of nominal compositions: Ti25Al25Nb(at.%), Ti23Al27Nb(at.%), and Ti12Al38Nb(at.%). Smaller cigar-shaped Ti25Al25Nb ingots were examined to understand the effect of rolling preheat treatment on microstructure. It was found that super-transus preheat treatment results in large prior BCC grains and surface edge cracking. For larger castings, forging and rolling procedures were carried out after heating the materials between 932–1000°C. These temperatures were below the BCC-transus temperature for Ti23Al27Nb and Ti25Al25Nb and above the transus for Ti12Al38Nb. This resulted in a signiﬁcantly larger grain size for the as-processed Ti12Al38Nb compared with the other two alloys. The Ti25Al25Nb alloy required the greatest forging and rolling loads, while the fully-BCC Ti12Al38Nb alloy exhibited the best workability and required the lowest forging and rolling loads. This was related to the alloys’ aluminum contents and O-phase volume fractions. Sub-transus processing of the near Ti2AlNb alloys proved to be a viable technique for obtaining homogeneous microstructures containing ﬁne O and BCC phases and lacking large prior BCC grains, which can be detrimental to the mechanical performance. © 2000 Elsevier Science S.A. All rights reserved.

Keywords: Titanium alloys; BCC phases; Orthorhombic phase

1. Introduction the thermomechanical processing has been based on previous methodologies developed for conventional a– b a Since the discovery of the orthorhombic (O) phase in titanium alloys and the intermetallic 2 titanium aTi25Al12.5Al(at.%)1 alloy by Banerjee et al. [1], aluminides. Early alloy development efforts focussed on titanium aluminides containing the O phase (based on extrusion, forging, or rolling operations on arc-melted ingots, with the primary intent being to characterize the Ti2AlNb) have been of interest for high-temperature structural applications, primarily because of their high equilibrium phases, phase transformations, and me- speciﬁc strength and stiffness as well as their creep and chanical behavior [5–8,10,11]. Later, Smith et al. [2–4] oxidation resistance. Recent results have shown that O used foil processing to examine the development of alloys offer major performance improvements over microstructure in Ti 22Al 23Nb and Rhodes et al. [12] commercial titanium alloys [2–10]. Like commercial studied microstructural evolution and crystallographic titanium alloys, the properties of O alloys depend texture in Ti 22Al 23Nb and Ti 22Al 27Nb sheet and strongly on the microstructure and therefore the pro- foil products during hot rolling, cold rolling, and subse- cessing. To date, the relationship between the process- quent heat treatment. More recently detailed studies of ing and microstructure of O alloys has been the development and control of microstructure during forging of a Ti22Al27Nb alloy [9] and hot pack investigated to a limited extent [1,5,6,8–16]. Much of rolling of a Ti22Al23Nb alloy have been performed [13]. These studies have focussed on a relatively narrow * Tel.: +1-410-5162876; fax: +1-410-5167316. a E-mail address: [email protected] (C.J. Boehlert) range of O alloy compositions, containing the O, 2 (ordered hexagonal close packed), and body-centered 1 All alloy compositions are given in atomic percent. cubic (BCC) phases, which are being considered for

metal matrix composite applications using foil-fiber-foil tions of the near Ti2AlNb alloys were Ti 25Al 25Nb processing. and Ti23Al27Nb. The initial portion of this study In this study, the processing-microstructure relation- involved examining the effect of processing parameters, ship for O alloys containing a wide range of composi- and in particular the rolling preheat temperature, on tions; namely Ti25Al25Nb, Ti23Al27Nb, and microstructure using smaller ingots prior to attempting Ti12Al38Nb alloys, was examined. Due primarily to large-scale deformation on larger castings. The term the high Nb content in such alloys, the two-phase used to describe the smaller ingots is ‘cigar melts’ O+BCC regime has a wider temperature range than because the dimensions, 150 mm in length and 30 mm a a the 2 +B2, 2 +B2+O, and O-phase fields. Hence diameter, are similar to that of a cigar. Three 300-g such alloys are termed ‘O+BCC’ alloys. The evolution cigar-melts were triple-melted using a vacuum induction of microstructure, from melted ingot to forged pancake melter at the Air Force Research Laboratory Materials to rolled sheet on the order of millimeters thick, was Directorate of Wright–Patterson Air Force Base, OH. examined. In attempt to avoid large prior-BCC grains, The heats were formulated using elemental Ti, Al, and which have shown to be detrimental to mechanical Nb according to the stoichiometric mixture Ti2AlNb behavior [17–19], conservative thermomechanical pro- and their measured compositions are listed in Table 1. cessing techniques comprising non-isothermal forging The larger castings consisted of 175–500 mm long and hot pack rolling at relatively low processing tem- cylinders of 75 mm diameter. The large, near Ti2AlNb peratures were chosen. In addition to further develop- ingots, whose compositions are listed in Table 2, were ing the understanding of microstructural evolution ‘induction-skull’ melted at Flowserve (formerly Duriron during processing of O+BCC alloys, this work de- Corp.), Dayton, OH. Note that for Ti23Al27Nb, the scribes how processing temperature affects the ability to Ti, Al, and Nb contents adhered well to the target control microstructural features, especially grain size, composition, while for Ti25Al25Nb, the measured which strongly influence the mechanical behavior. composition was close to Ti25Al23Nb. The large Ti12Al38Nb ingot, whose measured composition was close to Ti13Al39Nb (see Table 2), was vacuum arc 2. Experimental procedures melted at Pittsburgh Materials Technology Inc., Large, PA. Several samples, diamond cut from each material, 2.1. Materials and microstructural characterization were analyzed for their constituent elements. The Ti, Al, Nb, and Fe contents were analyzed by means of The studied alloys were grouped into two categories: solution X-ray fluorescence spectrometry and the data near Ti AlNb and Ti12Al38Nb. The target composi- 2 were obtained using a Kerex Corporation Model 770 Table 1 Delta Analyst. The amounts of nitrogen and oxygen were quantified using a Leco Corporation Model TC- Chemical analysis of the Ti2AlNb cigar-melted ingots and the corresponding as-processed sheetsa 136 oxygen/nitrogen analyzer. Chemical composition distribution between the different phases was measured Material Atomic percent Weight (ppm) using a Japan Electron Optics Ltd electron microprobe TiAl Nb N Fe H O analyzer (JEOL 733). Grain size (d) and phase volume fractions were determined quantitatively using NIH Ingot ABal 24.8 24.5 140 nana 250 image analysis software of digitized, high-contrast, Sheet ABal 25.4 24.2 110 350 na 280 back-scattered-detector (BSD) images taken using a Ingot BBal 26.6 23.5110 na 650 135 Leica 360 field-emission scanning electron microscope Sheet BBal 26.2 25.1 110 460 na 230 Ingot CBal 24.6 25.1 100 na na 290 (SEM). Transmission electron microscopy (TEM), per- Sheet C Bal 24.6 25.0 110 530 na 890 formed using a JEOL JEM-2000FX electron microscope, and X-ray diffraction (XRD) were used to a na, not available. confirm the presence of the different phases.

Table 2 2.2. Procedures Chemical analysis of the large ingots

Material Atomic percent Weight (ppm) 2.2.1. Forging and rolling procedures for the cigar-melted ingots TiAl Nb N Fe O Following melting, the cigar melts were cut by a wire electron discharge machine (EDM) to a rectangular Ti25Al25NbBal 24.7 23.3 150 290 930 Ti23Al27NbBal23.2 27.2 200 1100 1160 geometry, measuring 125×25×20 mm, and coated Ti12Al38Nb Bal 13.2 39.2 70 255 575 with high-temperature glass for lubrication and protection from the environment. They were then sealed in 6 120 C.J. Boehlert / Materials Science and Engineering A279 (2000) 118–129

Table 3

Rolling procedures and parameters for the Ti2AlNb cigar-melted ingots

Ingot A B C

Preform Forged (50%)Forged (50%) Forged (50%) Can dima (mm3) 140×60×12 122×66×12 150×66×12 Pre-rolling – — 1200°C/24 h Heat treatment Rolling-preheat 815C/1h/1000°C/0.25 h1040C/1h/1000°C/0.25 h 1040°C/1h/1000°C/0.25 h Interpass reheating 1000°C/2–3 min 1000°C/2–3 min 1000°C/2–3 min Intermediate anneal – 1060°C/0.5 h/1000°C/0.25 h 1060°C/.5 h/1000°C/0.25 h Reduction per pass 10%10% 10% Roll speed (m2 min−1) 2.3 2.3 2.3 Can dim (mm3) 412×64×3391×66×3 419×74×3 Sheet dimb(mm3) 381×46×2.7353×48×2.8 389×50×2.8

a After forging, the cans were weld repaired, reevacuated, and sealed for use in rolling. b All rolled sheets were reheated at 1000°C for 3 min and then cooled slowly in vermiculite (3°C min−1). mm thick stainless steel cans and unidirectionally was performed on a cold die and therefore after each forged according to a 2:1 ratio from 25 to 12.5 mm in pass the pancakes were reheated at 1000°C for 2–3 air at a rate of 150 mm min−1. Preheat treatment of the min. After finishing half of the rolling passes, each of ingots included an isothermal soak at 1050°C for 15 which resulted in a 10% reduction in thickness, pan- min followed by a 1000°C soak for 2 min prior to cakes B and C were annealed at 1060°C for 30 min and forging. After forging, the pancakes were slowly cooled then reheated at 1000°C for 15 min before continuing in vermiculite2. The cans were weld repaired, re-evacu- the remaining rolling passes. As a result of the rolling ated and sealed at room temperature (RT), then sub- operation, significantly longer workpieces were ob- jected to different preheat treatments prior to tained and the resulting sheets were approx. 3 mm subsequent rolling. The rolling operations, performed thick. Fig. 1 compares the size of a forged and forged- using a two-high laboratory mill, were chosen in order and-rolled cigar-melted ingot. to input more work into the forged pancakes and further homogenize the microstructures. Pack rolling was used to minimize the temperature transients that 2.2.2. Forging and rolling procedures for the larger occur during processing of thin sheet [13,20]. The ingots rolling procedures for each pancake (labeled A, B, and Based on the low-oxygen pick-up, the relatively low C) are given in Table 3. Three different heat-treatment degree of edge cracking (see Fig. 2), the fine-grained and rolling schemes were chosen to produce a range of microstructure (see Fig. 3a), and the balance of RT microstructures. The heat treatments consisted of low- tensile strength (1237 MPa) and elongation (5%) exhib- temperature (A), intermediate-temperature (B), and ited by sheet A [21], a similar rolling and forging high-temperature (C) exposures in air followed by a schedule was devised for the larger ingots. Forging sub-transus preheat stage in the O+BCC regime preforms, 60 mm in diameter and 150 mm tall, were (1000°C for 15 min) prior to rolling. Preheat treatment EDM cut from each of the larger castings and coated of pancake A comprised an initial sub-transus preheat- with high-temperature glass then sealed in 6 mm thick ing stage (815°C for 1 h) to dissolve most of the BCC stainless steel cans for protection from the environment. phase. The objective of this heat treatment and rolling The can assemblies, made up of both the workpiece and cycle was to produce a fine-grained O microstructure the outer can, were unidirectionally forged (3:1 ratio) to with a small volume fraction of BCC-phase particles. a final height of 50 mm in air at a rate of 150 mm Preheat treatment of pancake B consisted of an initial preheating stage near the BCC transus (1040°C for 1 h) to stabilize a larger BCC-phase volume fraction while limiting grain growth. Pancake C received a 24 h pre-rolling heat treatment at 1200°C, which is well above the BCC-transus temperature, in order to homogenize a large-grained, fully-BCC microstructure. This was followed by the same preheat treatment schedule as that of pancake B prior to rolling. The rolling Fig. 1. Size comparison of (a) forged and (b) forged and rolled 2 The estimated cooling rate was 3°C min−1. cigar-melted ingots. C.J. Boehlert / Materials Science and Engineering A279 (2000) 118–129 121

not undergo uniform deformation, no surface cracks were observed on these workpieces after the initial forging. This is discussed in the results and discussion section, which describes the added procedures used to process the Ti25Al25Nb workpiece. The Ti23Al27Nb and Ti12Al38Nb workpieces were then re-canned and forged to 25 mm in a direction perpendicular to that of the original under identical forging conditions. Similar to the first forging run, the can assemblies were uniformly deformed. Again the work pieces were removed and the sides of the pancakes were cut parallel to a width of 75 mm. After forging, the EDMed Ti23Al27Nb and Ti25Al25Nb pancakes were re-canned and isother- mally soaked at 815°C for 1 h followed by a 982°C soak for 15 min prior to rolling. The unidirectional rolling steps consisted of several passes on a cold die each after a soak at 982°C for 5 min. For the Ti12Al38Nb pancake, the soak temperature (932°C) was also identical to the forging temperature used for this alloy. The reduction per pass for each pancake was between 5–10%. The total reduction was approx. 60% and the total shear strain, calculated for both the Fig. 2. Low magnification photograph of the three as-rolled Ti AlNb 2 forging and rolling procedures, was on the order of sheets. three. The rolling loads were recorded for each pass on a stripchart and measured between 50–70 tons depend- ing on the reduction and the alloy. After the final pass, the sheets were reheated at the respective soaking temperatures for 3 min and then removed from the furnace and hung vertically for 1 min to promote creep straight- ening. They were then cooled in vermiculite. The final thickness of the sheets was approx. 12 mm and no further effort to produce thinner sheets or foils was made.

3. Results and discussion Fig. 3. As-rolled microstructures of sheets (a) A (b) B, and (c) C. These BSD SEM images were taken from the thickness section and 3.1. Microstructural e6olution the rolling direction is horizontal. 3.1.1. Cigar-melted ingots min−1. For the near Ti AlNb alloys, prior to forging 2 A qualitative assessment of the effect of the rolling the can assembly was heat treated at sub-transus tem- preheat treatment and the intermediate annealing steps peratures, 1000°C for 15 min followed by a 982°C soak is shown in Fig. 2, which depicts the three as-rolled for 2 min. For the Ti12Al38Nb alloy, the can assem- sheets fabricated from the cigar-melted ingots. Note bly was heat treated at 950°C for 15 min followed by a that the greatest amount of edge cracking was exhibited 932°C soak for 2 min prior to forging. Note that by the sample which underwent the most severe pre- although the forging temperature for Ti 12Al 38Nb rolling heat treatment (1200°C/24 h). Fig. 3a–c depicts was lower than that for the near Ti2AlNb alloys, it was the corresponding as-rolled sheet microstructures. For above the BCC-transus due to the decrease in BCC- each of the cigar-melted ingots the forging and rolling transus temperature with decreasing Al content [19,22]. procedures successfully broke down the O+BCC After forging, each pancake was cooled in vermiculite. platelet morphology, however some chemical banding, The Ti23Al27Nb and Ti12Al38Nb workpieces, indicated by the alternating light and dark layers in the which underwent uniform deformation, were removed BSD SEM images of Fig. 3a–c, remained. Table 1 lists from the can assemblies and EDM cut to a height of 75 the measured compositions of each of the ingots after mm. Unlike the Ti25Al25Nb workpieces, which did processing. The largest oxygen increase due to process- 122 C.J. Boehlert / Materials Science and Engineering A279 (2000) 118–129

second phase was found to be a combination of globu- lar/elongated particles within the grains as well as layers decorating the prior-BCC grain boundaries, see Fig. 3c. The prior-BCC grain sizes of sheets A and B, which did not undergo the 1200°C heat treatment, were more than an order of magnitude finer than that for sheet C. For sheet A, fine O- and BCC-phase particles were present throughout the microstructure. The as-processed sheet B microstructure was severely segregated and contained particles which were highly concentrated with Al and Nb (Ti32Al34Nb), see white particles of Fig. 3b. This segregation was determined to have been a result of the inhomogeneous starting ingot material [19,21], and it was also considered to be the reason why the large difference in Nb content was measured for the pre- and post-processed materials (see Table 1). Due to the fine, equiaxed microstructure, similar processing Fig. 4. BSD SEM image of the sheet A microstructure after a conditions, and in particular the rolling operations, as sub-transus heat treatment of 975°C for 100 h followed by water those used for sheet A, which was devoid of large quenching. prior-BCC grain boundaries and severe segregation, were chosen to process the larger-scale ingots. It is ing was measured for sheet C, which is expected to have noted that the chemical banding present in the as-pro- occurred during the 1200°C exposure for 24 h, and as a cessed sheet A was removed through sub-transus heat a result the 2-phase precipitated at near-surface loca- treatment, see Fig. 4, and a balance of RT tensile tions, see the bottom of Fig. 3c. Also due to this severe strength and elongation and elevated-temperature creep preheat treatment, the as-rolled microstructure of sheet resistance resulted [21]. C retained large elongated grains, which were on the order of 500 mm in length. The elongated grains in Fig. 3c suggest that a substantial amount of deformation 3.1.2. Larger ingots (approx. 4:1 thickness reduction) had been imposed The lower Al-containing ingots, Ti23Al27Nb and during rolling and that BCC-phase recrystallization did Ti12Al38Nb, were uniformly deformed around their not take place during deformation as large BCC grains centers during forging. The post-forged can assembly were present after the 1200°C/24 h heat treatment. The for Ti12Al38Nb is depicted in Fig. 5a and b. The

Fig. 5. The (a) top and (b) side views of the can assembly after the ﬁrst forging run for Ti12Al38Nb. C.J. Boehlert / Materials Science and Engineering A279 (2000) 118–129 123

Fig. 6. The (a) top and (b) side views of the extracted workpiece after the second forging for Ti12Al38Nb. circular cross-section of the workpiece can be clearly with respect to the forging direction. Photos of the seen in Fig. 5a, while a bulge in the stainless steel can is post-forged can assemblies for Ti25Al25Nb are de- present at its midpoint, see Fig. 5b. Fig. 6a and b depict picted in Fig. 9a and b. The circular shape of the top of the extracted Ti12Al38Nb workpiece after the second the severely displaced workpiece is easily recognized in forging run. The as-cast and forged pancake mi- Fig. 9a. The shearing resulted in a displacement of the crostructures for Ti12Al38Nb and Ti23Al27Nb are stainless steel can which bulged on opposite sides at the depicted in Fig. 7a and b and Fig. 8a and b, respec- top and bottom of the can assembly, see Fig. 9b. This tively, and the O and BCC phases were identified using behavior was reproducible as an identically configured both TEM and XRD. It is evident in Fig. 7a and b that workpiece and can exhibited similar behavior under the the BCC grain size was significantly reduced for same forging conditions. However, in this case, the Ti12Al38Nb. TEM investigations revealed that the BCC phase was disordered (designated as b), a result of the low Al content [19,22]. For Ti23Al27Nb, most of the prior-BCC grain boundaries were serrated during forging and although elongated O+BCC platelets remained, they were shorter and more ‘blockier’ than those in the as-cast material. Thus, the ingot microstructure had been broken down and the orientation of the platelets was more random than that of the as-cast microstructure. It appeared that neither the O nor BCC phase had been recrystallized and that the structure was merely a more highly wrought version of the as-melted ingot. Fig. 7. Comparison of the Ti12Al38Nb (a) as-cast and (b) forged Ti25Al25Nb did not undergo uniform deformation microstructures. The forging direction was vertical. Note the reduc- during forging and sheared at an approx. 45° angle tion in grain size after forging. 124 C.J. Boehlert / Materials Science and Engineering A279 (2000) 118–129

Fig. 8. Comparison of the Ti23Al27Nb (a) as-cast and (b) forged microstructures. The forging direction was vertical.

Fig. 9. The (a) top and (b) side views of the can assembly after the first forging run for Ti25Al25Nb. Note the non-uniform deformation. forging press was halted immediately after the initial shear deformation was observed, see Fig. 10. The onset of such instability and flow localization is typically observed for materials which exhibit high degrees of flow softening and/or low strain rate sensitivities [23]. High degrees of flow softening in conventional ab titanium alloys have several main sources including (i) forging at low temperatures and high strain rates at which the flow stress and deformation heating-induced softening are high or (ii) microstructural based softening such as occurs during the break-down of coarse- grain lamellar microstructure [23]. In this case the former source is more likely as both near Ti2AlNb ingots contained similar Widmanstatten microstructures, yet Ti25Al25Nb exhibited a greater compressive flow stress than Ti23Al27Nb, consistent with Fig. 10. The Ti25Al25Nb can assembly after an interrupted forging previous results on these materials [19]. Thus the dis- run. The forging direction was horizontal. Note the non-uniform placement rate of 150 mm min−1 or an approx. 0.017 deformation. C.J. Boehlert / Materials Science and Engineering A279 (2000) 118–129 125

depicted in Fig. 11a–c. Along with the shearing deformation, severe cracking occurred along the top surface of the workpiece, see Fig. 11a. Such shearing and cracking have been observed in isothermal forging of conventional a–b titanium alloys which also has been related to high rates of ﬂow softening and low values of the strain rate sensitivity index, m [23]. To prevent non-uniform deformation behavior, both the can assembly height and the reduction ratio were reduced prior to the second forging attempt. The slightly deformed workpiece, taken from the interrupted forging run (see Fig. 10), was EDM cut into two pieces, which were subsequently re-canned to forge in the same direction as that of the ﬁrst. The can assembly prior to forging was approximately 62 mm tall and the intended reduction ratio was 2:1. All other parameters, including forging speed, remained constant. The forging deformation was uniform for this run as depicted in Fig. 12a and b. Thus by simply reducing the can assembly height, uniform forging deformation was possible at 982°C. To acquire a similar amount of deformation for Ti25Al25Al as that for the other alloys, the Ti25Al25Nb workpieces were again removed, EDM cut, and re-canned in preparation for a third forging trial. The third forging step was performed in a direc-

Fig. 11. The (a) top (b) bottom and (c) side views of the sheared Ti25Al25Nb workpiece after the initial forging run. Note that the side depicted in (c) was EDM sliced after forging. s−1 compressive strain rate, which did not induce instability and ﬂow localization for Ti23Al27Nb, was too rapid for the stronger Ti25Al25Nb ingot at 982°C. This indicates the important inﬂuence of deformation heat-induced softening during forging. The extracted Fig. 12. The can assemblies of the two uniformly deformed Ti25Al25Nb workpiece after the initial forging run is Ti25Al25Nb pancakes after the second forging runs. 126 C.J. Boehlert / Materials Science and Engineering A279 (2000) 118–129

Fig. 13. Comparison of the Ti25Al25Nb (a) as-cast and (b) forged microstructures. The forging direction was vertical. tion perpendicular to the first two and the reduction tions of more than 5% were not achievable for ratio was again 2:1. The final height of the uniformly Ti25Al25Nb. The lowest rolling loads and best work- deformed can assembly was approx. 25 mm. Similar to ability were exhibited by Ti12Al38Nb, which was that of Ti23Al27Nb, the ingot microstructure was capable of single-pass rolling reductions of 10 percent. successfully broken down, see Fig. 13a and b. However, Fig. 14a and b depict the can assembly prior to rolling in this case the second-phase BCC particles were more and the post-rolled workpiece for Ti12Al38Nb. Note spherical than those for Ti23Al27Nb. Also note the the substantial lengthening which resulted from rolling. lower BCC-phase volume fraction for Ti25Al25Nb Fig. 15a and b depict the can assembly prior to rolling (Vf 0.05) compared to that for Ti 23Al 27Nb (Vf and the post-rolled workpiece for Ti23Al27Nb, while 0.30), compare Fig. 8 and Fig. 13, which is a result of Fig. 16a and b depict the post-rolled can assembly and the phase equilibrium [19,22]. Thus, the Ti25Al25Nb the extracted workpiece for Ti25Al25Nb. Unlike the material was workable at 982°C, though the ingot rolling procedures for the cigar-melted ingots, the height needed to be reduced from that used for the sheets from the larger ingots exhibited almost no edge Ti 23Al 27Nb and Ti 12Al 38Nb materials. It should cracking. Thus, O+BCC alloys are quite amenable to be noted that the applied loads for the initial forgings intermediate-temperature rolling with the microstruc- were greatest for Ti25Al25Nb and smallest for Ti12Al38Nb. The reason for the poorer workability of the Ti25Al25Nb ingot is believed to be due to the higher Al content and the corresponding greater O- phase volume fraction. The poor workability of alloys containing high Al concentrations has been reported previously [24], and microstructures containing greater volume fractions of the undecomposed BCC phase exhibit lower isothermal compressive flow stresses [15], which explains why the Ti23Al27Nb (which contained an intermediate BCC phase volume fraction) exhibited loads intermediate to those for Ti25Al25Nb and Ti12Al38Nb. Ti12Al38Nb, which was fully-b at the forging and rolling temperature, was the most amenable to forging and this is related to the low flow stress and excellent ductility of the b phase. On the other hand, the BCC phase of Ti25Al25Nb and Ti23Al27Nb was ordered (B2) at the forging and rolling temperatures [19,22] and this is expected to have influenced the applied loads as well. Similar to the forging procedures, the largest rolling loads and poorest workability during rolling were ex- Fig. 14. The (a) can assembly prior to rolling and the (b) extracted hibited by Ti25Al25Nb. Single-pass rolling reduc- workpiece after rolling for Ti12Al38Nb. C.J. Boehlert / Materials Science and Engineering A279 (2000) 118–129 127

exception that the O and BCC phases were elongated in the rolling direction. The rolled Ti23Al27Nb microstructure also contained elongated O+BCC phases, Fig. 17b, yet they were less elongated, or more oval shaped, than those for the forged material, see Fig. 8b. The development of this oval-like morphology, rather than the platelet morphology, may be attributed to the large amount of hot work during forging and rolling and lower amounts of growth due to sub-transus processing temperatures. In addition, recrystallization of the matrix phase was not apparent. Thus, due to the sub-transus rolling and forging steps, the platelet-like

morphology of the near Ti2AlNb ingots evolved to a more equiaxed microstructure. The largest average

grain dimension of the near Ti2AlNb as-processed microstructures was approximately 4 mm and a greater BCC-phase volume fraction was evident for Ti23Al27Nb compared to Ti25Al25Nb. Overall, the degree of homogeneity in the microstructure was high after rolling for these sheets. Further rolling to thinner sheet or foil, through either hot rolling or cold Fig. 15. The (a) can assembly prior to rolling and the (b) extracted rolling [12,14], would increase homogeneity en route to workpiece after rolling for Ti 23Al 27Nb. quality foils used in foil-ﬁber-foil processing. It is noted that both sub-transus and super-transus heat treatments of the as-processed sheets resulted in homogeneous microstructures and a detailed description of the phase evolution, which includes the temperature ranges for the different phase regimes, is provided in Ref. [22]. A detailed description of the creep and tensile behavior of these alloys is provided in Ref. [19] and [25]. For the super-transus processed Ti12Al38Nb, the as-rolled microstructure, see Fig. 17c, contained fully-b grains elongated in the rolling direction. However, the overall grain dimensions were over an order of magnitude ﬁner than those of the as-cast microstructure, see Fig. 7a, and the largest average grain dimension was approx. 30 mm. The highly elongated b grains suggested an absence of recrystallization. Thus by performing forging and rolling at a relatively low temperature (932°C) within the single-phase b region (the b-transus

Fig. 16. The post-rolled (a) can assembly and (b) extracted workpiece for Ti25Al25Nb. tures containing larger BCC-phase volume fractions providing better workability, which is consistent with the observations of Rhodes et al. [12]. Low magniﬁca- tion BSD SEM images of the three as-rolled sheet Fig. 17. As-rolled microstructures of the (a) Ti25Al25Nb (b) microstructures are depicted in Fig. 17a–c. The Ti23Al27Nb, and (c) Ti12Al38Nb sheets. These BSD SEM im- Ti25Al25Nb sheet microstructure, Fig. 17a, was quite ages were taken from the thickness section and the rolling direction is similar to the forged material, see Fig. 13b, with the horizontal. 128 C.J. Boehlert / Materials Science and Engineering A279 (2000) 118–129 temperature was determined to be 800°C [22]), followed for equal deformation. Ti25Al25Nb required a by a short time hold at the processing temperature, the reduced forging height than Ti23Al27Nb for uni- as-cast grain size was signiﬁcantly reduced while main- form deformation, suggesting a difference in the taining a homogeneous single-phase microstructure. deformation heat-induced softening behavior with increasing Al content and corresponding higher O- phase volume fraction. 4. Summary and conclusions

The microstructural evolution during hot upset forg- Acknowledgements ing and hot pack rolling was investigated for O+BCC TiAlNb alloys. The alloys examined were This research was performed at the Wright-Patterson Ti25Al25Nb, Ti23Al27Nb, and Ti12Al38Nb Air Force Research Laboratory Materials and Manu- nominally. The former two were grouped as near facturing Directorate under Air Force contracts

Ti2AlNb alloys. Smaller ingots were used to understand F33615-91-C-5663 and F33615-C-96-5258 to UES, Inc. the effect of rolling preheat treatment on as-rolled The author is especially grateful to Dr V. Seetharaman microstructure for Ti25Al25Nb. The largest grain for technical guidance and Drs B.S. Majumdar, D.B. a size, oxygen pick-up, surface 2 precipitation, and sur- Miracle, and S.L. Semiatin for helpful discussions. The face edge cracking was exhibited for the super-transus assistance of J. and T. Brown, T. Jones, and T. Goff of preheated sheet. Each of the larger castings was pro- UES., Inc. in conducting the cigar melting, forging, and cessed to sheet at temperatures between 932–1000°C. rolling experiments is gratefully acknowledged. The Hot forging procedures initiated the breakdown of the author would also like to acknowledge the support large prior BCC grains of the ingots, while hot pack received from Johns Hopkins University during the rolling further reduced the grain size. The writing of this manuscript. Ti12Al38Nb alloy was the easiest to deform, which is expected to be a result of two main factors: (i) the flow References stress of the near Ti2AlNb O+BCC alloys is significantly higher than that of the fully-b Ti12Al38Nb; (ii) the fully-b Ti12Al38Nb microstructure exhibits [1] D. Banerjee, A.K. Gogia, T.K. Nandy, V.A. Joshi:, Acta Metall. 36 (4) (1988) 871–882. excellent ductility [19,26] and was processed above the [2] P.R. Smith, J.A. Graves, C.G. Rhodes, Metall. Trans. 25A b transus, while the near Ti2AlNb alloys were processed (1994) 1267–1283. at sub-transus temperatures. [3] P.R. Smith, W.J. Porter, W.J. Kralik, J.A. Graves, WL-TR-95- The following lists the conclusions of this work. 4068, Wright Patterson Air Force Base, OH, 1994, pp. 371–85. 1. The achievability of desired microstructures is [4] P.R. Smith, W.J. Porter, W.J. Kralik, J.A. Graves, Metal matrix composites, in: A. Poursartip, K.N. Street (Eds.), Proceedings of strongly dependent on the processing and heat-treat- the Tenth International Conference on Composite Materials, ment schedules. Homogeneous, fine-grained (d4 vol. 2, Woodhead, Cambridge, UK, 1995, pp. 731–738. m m) microstructures of near Ti2AlNb alloys were [5] R.G. Rowe, D. Banerjee, K. Muraleedharan, M. Larsen, E.L. produced through sub-transus processing, which re- Hall, D.G. Konitzer, A.P. Woodfield, in: F.H. Froes, I. Caplan sulted in the break down of the prior BCC grain (Eds.) Titanium ‘92 Science and Technology, The Minerals, Metals, and Materials Society, 1993, pp. 1259–66. boundaries and platelet-like morphology. Mi- [6] R.G. Rowe, P.A. Siemers, M. Larsen, Advances in the Process- crostructures containing large prior BCC grains ing, Synthesis, Characteristics, and Applications of Aerospace were produced when super-transus processing heat- Metal Based Materials, Proceedings Third International SAMPE treatment temperatures were used. It is concluded Metals and Metals Processing Conference, 1992. that fine-grained microstructures are possible only [7] R.G. Rowe, Physical Metallurgy Laboratory, GE Re- portc93CRD030, 1993. when work is performed below the transus. [8] C.M. Austin, J.R. Dobbs, H.L. Fraser, D.G. Konitzer, D.J. 2. Super-transus forging and rolling of Ti12Al38Nb Miller, M.J. Parks, J.C. Schaeffe, J.W. Sears, Rapidly Solidified produced intermediate grain sized (d30 mm) fully- Oxidation Resistant Niobium Base Alloys, WL-TR-93-4059, GE b microstructures. Due to the excellent ductility and Aircraft Engines, Cincinnati, OH, 1992. lower yield stress of the low-Al b phase, this mate- [9] A.P. Woodfield, Progress Report No. 5, General Electric Air- craft Engines, Cincinnati, OH, 1996. rial exhibited better workability than the O+BCC [10] J.C. Chesnutt, R.A. Amato, C.M. Austin, R.L. Fleischer, near Ti2AlNb alloys. M.F.X. Gigliotti, D.A. Hardwick, S.C. Huang, D.G. Konitzer, 3. The BCC phase does not recrystallize during sub- M.M. Lee, P.L. Martin, C.G. Rhodes, R.G. Rowe, G.K. Scarr, transus hot forging and hot rolling. The absence of D.S. Shih, P.A. Zomcik, Very High Temperature Titanium-Base recrystallization leads to microstructures with elon- Materials Research, WL-TR-91-4070, GE Aircraft Engines, Cincinnati, OH, 1993. gated BCC grains. [11] D. Banerjee, T.K. Nandy, A.K. Gogia, K. Muraleedharan, 4. Greater Al contents and O-phase volume fractions Titanium ‘88 Science and Technology, The Minerals, Metals, lead to greater forging and rolling loads necessary and Materials Society, 1989, 1091–96. C.J. Boehlert / Materials Science and Engineering A279 (2000) 118–129 129

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