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  • Acta metall, mater. Vol. 40, No. 11, pp. 3035-3049, 1992 0956-7151/92 $5.00 + 0.00 Printed in Great Britain. All rights reserved Copyright 1992 Pergamon Press Ltd


    L. M. HSIUNG, W. CAI and H. N. G. WADLEY Department of Materials Science and Engineering, University of Virginia, Charlottesville,

    VA 22903-2442, U.S.A.

    (Received 6 April 1992)

    Abstract--Isothermal phase transformations of a rapidly solidified Ti-24AI-I 1Nb alloy at temperatures below 900C have been studied using electron microscopy (SEM and TEM) and X-ray diffraction (XRD) methods. A supercooled fl phase was found to retain in the alloy as a result of rapid cooling. Subsequent isothermal treatments resulted in a fl ---+ ~2 transition which took place through a complex sequence of phase transitions. Initially it proceeded by ordering transitions fl (disordered b.c.c.)--, B2 (ordered b.c.c.) ---} T (ordered tetragonal) and was then followed by T ~ O (ordered orthorhombic) ---} ~2 (ordered hexagonal) transitions. Plate-like O and a2 phases were found to form at the early stages of both the T ~ O and O---~a 2 transitions. The T/O interface (habit) plane was determined to be {223}r, and the O/~ 2 interface (habit) plane was determined to be {lI0}o. Both the T----} O and O---} ~2 transitions can be explained by a shape deformation mechanism involving a homogeneous lattice distortion and a rigid-body rotation. In addition to the shape deformation, diffusion of Nb away from the O/~ 2 interface also is needed for the O----, ~2 transition to progress.

    R~um~--On 6tudie les transformations de phase isothermes d'un alliage Ti-24Al-11Nb, rapidement solidfi6 fi des temperatures infrrieures /t 900C, par microscopie 61ectronique en transmission et en balayage (MET et MEB) et par diffraction des rayons X. On trouve une phase fl surfondue qui reste dans I'alliage par suite du refroidissement rapide. Des traitements isothermes ultrrieurs provoquent la transition fl ~ ~2 qui se produit fi la suite d'une srquence complexe de transitions de phases: Au drbut, il se produit des transformations d'ordre fl (c.c. drsordonnr)---}B2 (c.c. ordonnr)---}T (quadratique ordonnr), suivies par les transitions T---~O (orthohombique ordonnr)---}~2 (hexagonal ordonnr). On trouve que les phases Oet ~2, en forme de plaquettes, se forment aux premiers stades des transitions T---}O et O--}a 2. Le plan d'interface (plan d'accolement) T/O est {223}r et le plan d'interface O/~ 2 est { 1T0}o. Les deux transitions T--*O et O---}~ 2 peuvent s'expliquer par un m~canisme de drformation de forme qui implique une distorsion homogrne et une rotation rigide du rrseau. En plus de la drformation de forme, la diffusion du Nb hors le l'interface O/~2 est aussi n~.essaire pour que la transformation O---~a2 s'effectue.

    Z~-,ammenfassung--Die isothermen Phasenumwandlungen in der rasch erstarrten Legierung Ti-24AI- 11Nb werden bei Temperaturen unterhalb von 900C elektronenmikroskopisch (Raster- und Durch- strahlungselektronenmikroskopie) und mit Rfntgenbeugungsmethoden untersucht. Als Ergebnis der raschen Erstarrung bleibt eine unterkfihlte fl-Phase in der Legierung zuriick. Nachfolgende isotherme W/irmebehandlung fiihrt zum ~bergang fl---}~2, der fiber eine komplexe Folge von Phaseniiberg/ingen l~iuft: zu Anfang l/iuft sic fiber Ordnungsumwandlungen fl (entordnet k.r.z.)---}B2 (geordnet k.r.z.)--+T (geordnet tetragonal), gefolgt von T---}O (geordnet orthorhombisch)---,a 2 (geordnet hexagonal). Platten- ertige O- und ~2-Phasen bilden sich im frfihen Stadium der Umwandlungen T---}O und O---~a 2. Die T/O-Grenzflfichen-(Habit)-Ebene wird zu {223} bestimmt, die O/~2-Ebene zu {li0}o. Sowohl der Obergang T--,O, wie auch O---*~ 2 krnnen mit einem Form-Deformationsmechanismus, der eine homogene Gitterverzerrung und eine Festkrrperrotation einschlieBt, erkl~irt werden. Zus/itzlich zur Formdeformation wird die Diffusion von Nb weg von der O/a2-Grenzfl/iche ffir den l~bergang O---~:c 2 benrtigt.


    Ti3Al-base aluminide alloys are candidate materials for high temperature structural applications because of their low density, good elevated temperature strength and creep resistance. Advanced intermetallic- matrix composites (IMCs) that use the Ti3AI aluminide alloys as a matrix can offer even better

    performance in stiffness and strength over their monolithic counterparts [1]. Ti3A1 aluminide alloys, however, are difficult to be fabricated into structural components by conventional processing methods due to limited ductility and low toughness, To overcome this, researchers are exploring the use of inductive- coupled plasma-deposition (ICPD) [1, 2] to produce composite unidirectional monotapes followed by hot


  • 3036 HSIUNG eta/.: RAPIDLY SOLIDFIED Ti-24Al-11Nb

    isostatic pressing (HIPing) or vacuum hot pressing (VHPing) to make a multilayer net shape component. Before one can control and optimize the matrix microstructures (and thus the IMC's mechanical properties) by controlling the deposition and subse- quent consolidation processes, a full understanding of time-temperature-transformation (TTT) as well as isothermal transformation behavior of the Ti a Al-base aluminide alloys is needed. The TTT behavior of different Ti3Al-base aluminide alloys have been investigated by Weykamp et al. [3], Peters and Bassi [4], Djanarthany et al. [5], and Hsiung et al. [6]. Nevertheless, inconsistency can still be found among their results in terms of transformation sequence. In addition, little has been mentioned about mechanisms, and a systematic study of the TTT behavior of the TiaAl-base aluminide system has yet to be reported.

    A candidate Ti3Al-base aluminide system used as the matrix material for fabricating IMCs is Ti3A1 with an addition of Nb. Nb is added to stabilize the fl phase and improve the deformability of the Ti3A1- base aluminide at lower temperatures [7]. Studies of microstructure in quenched (or rapidly solidified) and aged Ti 3 A1 + Nb alloys have revealed the existence of numerous phases [8-14]. These include fl (disordered b.c.c.), B2 (ordered b.c.c.), to-like (ordered h.c.p.),

    T (ordered tetragonal), O (ordered orthorhombic) and a2 (ordered h.c.p.) phases. Three different modes of isothermal phase transitions have been recognized in the Ti3A1 + Nb system [4, 6]: (I) fl --, ~2 + fl, (II) f l-- , a2(+O) and (III) fl--* to. For the case of Ti -24Al- l lNb [6], mode (I) was found to occur between 900 and 1150C, mode (II) below 900C, and mode (III) was found to be a competition reaction to the mode (II) reaction at temperatures below 650C. We report here on a detailed study of the mode (II) reaction. The effort has focused upon elucidating the mechanisms of the mode (II) reaction.


    A rapidly-solidified Ti3AI+ Nb aluminide alloy with a nominal composition of Ti -24at .%Al- l l a t .%Nb (Ti-14wt%A1-21 wt%Nb) was chosen for this study. Two forms of this alloy were used: powder and foil. The powder was produced from ingot via the plasma rotary electrode process (PREP) by Nuclear Metals Inc., Concord, Mass. The foil was produced from powder of identical composition via the inductively coupled plasma deposition (ICPD) process at GE Aircraft Engines, Lynn, Mass. [2]. During the ICPD process, powder was melted by


    P R E P P o w d e r

    ( b ) X R D

    i | I i , , 3 0



    2 0 0

    I I i ~ I I I ! I | I ! i I I I a

    ,40 5 0 6 0 7 0

    2 #

    Fig. I. (a) SEM micrograph showing microstructure of the as-prepared PREP powder, (b) XRD pattern (Cu K~) showing the existence of the B phase in the as-prepared powder.

  • HSIUNG et al.: RAPIDLY SOLIDFIED Ti-24AI-I 1Nb 3037

    passing it through a plasma. The molten droplets were then immediately deposited onto a preheated mandrel inside a vacuum chamber where they were rapidly quenched to a solid state. This mandrel was preheated to ~ 800C and held at this temperature during deposition. Thus, during the deposition process the alloy first deposited was cooled to ~800C, and was held at this temperature until completion of the deposition, and was then cooled to room temperature. The total deposition period lasted approximately 1 h. Therefore the alloy first deposited has already been preaged.

    Both the powder and foil were examined in their as- prepared and aged status using X-ray diffractometry (XRD), electron microscopy (SEM and TEM), and energy dispersive X-ray spectrometry (EDS). Prior to aging, specimens were wrapped with tantalum foils and encapsulated in cleaned and evacuated quartz ampoules. Aging was performed for various times in the temperature range 450 to 900C. Following the heat treatment, Vickers hardness of PREP powders was measured using a MICROMET microhardness

    B XRD i~ ICPD Foil: As-sprayed

    I (a) Side A B

    o ~ o B

    o 20 I

    O0 I ~0 (b)Center =( o o l

    ~ 0 0 ; k 042 0 431

    ~12~ 20 ":' (c) Side B


    o~' {9. a, a 2 a~ I "= I l l =o.= ==..o =o'=1

    =e 4o so do "Io

    2e Fig. 2. XRD pattern (Cu K=) generated from the as-sprayed ICPD foil, (a) side A (last deposited), (b) center portion and

    (c) side B (early deposited).

    indentor. TEM specimens were prepared from both the as-prepared and aged ICPD foil. Microstructures were examined using a Philips 400T transmission electron microscope. Selected-area electron diffraction (SAD), microdiffraction (MD) and convergent-beam electron diffraction (CBED) methods were applied to identify and distinguish the different phases.

    3. RESULTS

    3.1. As-processed microstructures

    3.1.1. PREP powder. The microstructure of the as- prepared powder was studied using XRD and SEM. The results are shown in F