Academia.eduAcademia.edu

Figure 13 – uploaded by Dan Eliezer

See full PDFdownloadDownload figure
Figure 14 Internal transverse matrix cracking in a Ti-24Al- 11Nb/SCS-6 composite experiencing 500 cycles in the temperature range of 150 to 815°C [263]. Figure 13 High-magnification micrograph of the reaction zone. Diffraction patterns on left are from the (Ti, Nb); AIC phase, which are seen as faulted grains in the micrograph. Patterns on the right are from the (Ti, Nb, Al);Si, particles, one of which is indicated with an arrow [252].

Figure 14 Internal transverse matrix cracking in a Ti-24Al- 11Nb/SCS-6 composite experiencing 500 cycles in the temperature range of 150 to 815°C [263]. Figure 13 High-magnification micrograph of the reaction zone. Diffraction patterns on left are from the (Ti, Nb); AIC phase, which are seen as faulted grains in the micrograph. Patterns on the right are from the (Ti, Nb, Al);Si, particles, one of which is indicated with an arrow [252].

Related Figures (29)

TABLE I Melting points and densities for aluminides [34]
TABLE I Melting points and densities for aluminides [34]
Figure 1 Comparison of the creep behaviour (0.2% creep strength) of terminal and intermetallic titanium-based alloys [9]. 1 lbin~? = 703 kgm7?.
Figure 1 Comparison of the creep behaviour (0.2% creep strength) of terminal and intermetallic titanium-based alloys [9]. 1 lbin~? = 703 kgm7?.
TABLE ITI Possible aircraft engine material mix in 2010 [34]
TABLE ITI Possible aircraft engine material mix in 2010 [34]
Figure 2 Material mix in a current advanced turbofan engine [34]. °C = (CF — 32)/1.8.
Figure 2 Material mix in a current advanced turbofan engine [34]. °C = (CF — 32)/1.8.
TABLE III Properties of titanium aluminides, titanium-base conventional alloys and superalloys [28-35] * Throughout the text all compositions are expressed in at % unless otherwise stated.
TABLE III Properties of titanium aluminides, titanium-base conventional alloys and superalloys [28-35] * Throughout the text all compositions are expressed in at % unless otherwise stated.
Figure 3 Schematic binary Ti-Al phase diagram [66]. Care should be exercised in evaluating early work because of variations in composition, processing, microstructure, and testing procedures [87]. Of par- ticular importance is oxygen level, which raises the ductile-to-brittle transition temperature of titanium alloys [87]. Reduction of oxygen content from 0.186 to 0.081 wt % for the Ti-24AI-11Nb alloy results in an increase in tensile elongation from 1.1% up to 4.8% Additions of B-stabilizing elements such as Nb, Mo, V and Ta enhance room-temperature ductility, and much work is in progress to establish the constitution of ternary Ti-AI-M (M = Nb, V, Mo, Ta, Mn, etc.) systems [68-74], with emphasis on Ti-AI-Nb [73, 74]. In Ti-Al-Nb alloys containing 5 to 17 at % Nb a ductile ordered b.c.c. phase (B2) forms on quen- ching from high temperatures [75]; this also occurs in alloys containing other B-stabilizing elements [76, 77]. Additionally, an equilibrium ordered orthorhombic phase is detected in Ti-AI-Nb alloys with an appro- ximate stoichiometry of Ti,AINb [78]. Alloys quen- ched from the high-temperature B phase (both in the binary and the ternary alloys) undergo a variety of martensitic transformations. Similarly, alloys contain- ing the B2 phase form the w and @-related phases on
Figure 3 Schematic binary Ti-Al phase diagram [66]. Care should be exercised in evaluating early work because of variations in composition, processing, microstructure, and testing procedures [87]. Of par- ticular importance is oxygen level, which raises the ductile-to-brittle transition temperature of titanium alloys [87]. Reduction of oxygen content from 0.186 to 0.081 wt % for the Ti-24AI-11Nb alloy results in an increase in tensile elongation from 1.1% up to 4.8% Additions of B-stabilizing elements such as Nb, Mo, V and Ta enhance room-temperature ductility, and much work is in progress to establish the constitution of ternary Ti-AI-M (M = Nb, V, Mo, Ta, Mn, etc.) systems [68-74], with emphasis on Ti-AI-Nb [73, 74]. In Ti-Al-Nb alloys containing 5 to 17 at % Nb a ductile ordered b.c.c. phase (B2) forms on quen- ching from high temperatures [75]; this also occurs in alloys containing other B-stabilizing elements [76, 77]. Additionally, an equilibrium ordered orthorhombic phase is detected in Ti-AI-Nb alloys with an appro- ximate stoichiometry of Ti,AINb [78]. Alloys quen- ched from the high-temperature B phase (both in the binary and the ternary alloys) undergo a variety of martensitic transformations. Similarly, alloys contain- ing the B2 phase form the w and @-related phases on
* Observed only in ternary Ti-AI-M alloys where M is a B-stabilizing elemen > Low-temperature equilibrium phase in Ti-AI-Nb alloys. * Transitional phases formed during decomposition of the B2 phase. TABLE IV Crystal structure data of the stable and metastable phases in titanium-rich Ti-Al and Ti-AI-Nb alloys
* Observed only in ternary Ti-AI-M alloys where M is a B-stabilizing elemen > Low-temperature equilibrium phase in Ti-AI-Nb alloys. * Transitional phases formed during decomposition of the B2 phase. TABLE IV Crystal structure data of the stable and metastable phases in titanium-rich Ti-Al and Ti-AI-Nb alloys
Figure 4 Microstructures developed in a Ti-25Al-7.5Nb alloy by transforming from the B phase: (a) equiaxed microstructure with a high volume fraction of primary «, phase, (b)low volume fraction of primary «, phase in a Widmanstatten transformed matrix, (c) coarse Widmanstatten transformed microstructure and (d) fine Widmanstatten transformed microstructure [82].
Figure 4 Microstructures developed in a Ti-25Al-7.5Nb alloy by transforming from the B phase: (a) equiaxed microstructure with a high volume fraction of primary «, phase, (b)low volume fraction of primary «, phase in a Widmanstatten transformed matrix, (c) coarse Widmanstatten transformed microstructure and (d) fine Widmanstatten transformed microstructure [82].
Figure 5 (a) Crystal structure of Ti;Al («,) phase and (b) possible slip planes and slip vectors in the structure [31]
Figure 5 (a) Crystal structure of Ti;Al («,) phase and (b) possible slip planes and slip vectors in the structure [31]
* Hours at 650°C/380 MPa. > Specially processed [112]. ° Estimated from microhardness [19]. TABLE V Typical room-temperature mechanical properties and elevated creep rupture life of a, [19, 31, 32, 82, 86, 112]
* Hours at 650°C/380 MPa. > Specially processed [112]. ° Estimated from microhardness [19]. TABLE V Typical room-temperature mechanical properties and elevated creep rupture life of a, [19, 31, 32, 82, 86, 112]
Figure 6 Variety of microstructures in y-TiAl alloys (unpublished work by C. Suryanarayana at WPAFB, Ohio, 1990) In two-phase alloys, the ratio of lamellar to equiaxed y must also be controlled [29—33, 119, 120] along with their distribution [124]. A lamellar volume fraction of about 30% gives a reasonable combination of properties with good high-temperature creep resist- ance and acceptable tensile strength and usable ductil- ity [29]. The optimum structure for enhanced ductility has been suggested to consist of adjacent regions of lamellar and equiaxed morphologies [33, 124]. TMP refines the microstructure when processing is con- ducted such that both the « and y grains are re- crystallized in the (« + y) phase field. Grain morpho- The ordered nature of TiAl (successive (002) planes are made up entirely of Ti and Al atoms) and its slight tetragonality. have important implications for the available Burgers vectors and the different stacking faults produced by partial dislocations. The
Figure 6 Variety of microstructures in y-TiAl alloys (unpublished work by C. Suryanarayana at WPAFB, Ohio, 1990) In two-phase alloys, the ratio of lamellar to equiaxed y must also be controlled [29—33, 119, 120] along with their distribution [124]. A lamellar volume fraction of about 30% gives a reasonable combination of properties with good high-temperature creep resist- ance and acceptable tensile strength and usable ductil- ity [29]. The optimum structure for enhanced ductility has been suggested to consist of adjacent regions of lamellar and equiaxed morphologies [33, 124]. TMP refines the microstructure when processing is con- ducted such that both the « and y grains are re- crystallized in the (« + y) phase field. Grain morpho- The ordered nature of TiAl (successive (002) planes are made up entirely of Ti and Al atoms) and its slight tetragonality. have important implications for the available Burgers vectors and the different stacking faults produced by partial dislocations. The
*Not Thompson notation. these produce complex stacking faults (CSF) and anti- phase boundary (APB) structures. Some typical ex- amples are as follows: Thompson tetrahedron, modified for the Ll) struc- ture, can be used to describe the Burgers vectors and he dissociation reactions of the dislocations (Fig. 7b). n this diagram, A, B, P and Q are the vertices of the etrahedron; «, B, y and 6 represent the centres of the faces opposite to the vertices A, B, P and Q, respect- vely. Any displacement vector connecting either A or B with P or Q leads to chemical disordering and herefore it is not likely that these dislocations will exist. Perfect dislocations are defined by two consecut- ive capital letters AB, PQ, etc; partial dislocations are described by mixed letters of one capital letter and one Greek letter, e.g. PB, BA, etc; and superdislocations are designated by mixed letter(s) and number (larger than unity), e.g. 2PB, 3BQ, etc. The Burgers vectors and magnitudes of the possible perfect dislocations in TiAl are presented in Table VI [131, 133, 134]. Movement of any of these dislocations does not destroy ordering. ~~. .
*Not Thompson notation. these produce complex stacking faults (CSF) and anti- phase boundary (APB) structures. Some typical ex- amples are as follows: Thompson tetrahedron, modified for the Ll) struc- ture, can be used to describe the Burgers vectors and he dissociation reactions of the dislocations (Fig. 7b). n this diagram, A, B, P and Q are the vertices of the etrahedron; «, B, y and 6 represent the centres of the faces opposite to the vertices A, B, P and Q, respect- vely. Any displacement vector connecting either A or B with P or Q leads to chemical disordering and herefore it is not likely that these dislocations will exist. Perfect dislocations are defined by two consecut- ive capital letters AB, PQ, etc; partial dislocations are described by mixed letters of one capital letter and one Greek letter, e.g. PB, BA, etc; and superdislocations are designated by mixed letter(s) and number (larger than unity), e.g. 2PB, 3BQ, etc. The Burgers vectors and magnitudes of the possible perfect dislocations in TiAl are presented in Table VI [131, 133, 134]. Movement of any of these dislocations does not destroy ordering. ~~. .
Figure 7 (a) Crystal structure and Burgers vectors in the TiAl (y) phase and (b) modified Thompson tetrahedron for the L1, structure [140° TABLE VI Magnitudes of the possible Burgers vectors in TiAl in ascending order; TiAl (with c/a = 1.02) is regarded as cubic (afte (131, 134])
Figure 7 (a) Crystal structure and Burgers vectors in the TiAl (y) phase and (b) modified Thompson tetrahedron for the L1, structure [140° TABLE VI Magnitudes of the possible Burgers vectors in TiAl in ascending order; TiAl (with c/a = 1.02) is regarded as cubic (afte (131, 134])
Figure 8 (a) Turbine engine compressor rotor produced by HIP of x, PREP powder and (b) machining preform produced by HIP of gas- atomized y powder using the ceramic mould process [180].
Figure 8 (a) Turbine engine compressor rotor produced by HIP of x, PREP powder and (b) machining preform produced by HIP of gas- atomized y powder using the ceramic mould process [180].
Figure 9 Er,O, dispersion in a rapidly solidified Ti,Al-2 wt %. Er alloy HIPed at 850°C [196].
Figure 9 Er,O, dispersion in a rapidly solidified Ti,Al-2 wt %. Er alloy HIPed at 850°C [196].
Figure 1] Comparison of o/p values of (@) SiC/Ti,Al + Nb (40 vol % SiC), (O) Ti;AJ] + Nb and (shaded) wrought René 41 and Hastelloy X superalloys for a range of temperatures [238]. Figure 10 Tensile strength versus temperature for (CJ) SiC fibre (0.014 cm), (O) Ti,Al + Nb, (@) SiC/Ti,Al + Nb (40 vol % fibre), and (A) ROM prediction (40 vol % fibre) [238].
Figure 1] Comparison of o/p values of (@) SiC/Ti,Al + Nb (40 vol % SiC), (O) Ti;AJ] + Nb and (shaded) wrought René 41 and Hastelloy X superalloys for a range of temperatures [238]. Figure 10 Tensile strength versus temperature for (CJ) SiC fibre (0.014 cm), (O) Ti,Al + Nb, (@) SiC/Ti,Al + Nb (40 vol % fibre), and (A) ROM prediction (40 vol % fibre) [238].
The advantages to be gained by embedding ceramic fibres in a titanium matrix have received considerable attention [231-237]. The enhancement in longitudinal tensile strength which is theoretically predicted from rule-of-mixtures (ROM) calculations, and has actually been achieved over a substantial temperature regime,
The advantages to be gained by embedding ceramic fibres in a titanium matrix have received considerable attention [231-237]. The enhancement in longitudinal tensile strength which is theoretically predicted from rule-of-mixtures (ROM) calculations, and has actually been achieved over a substantial temperature regime,
Figure 12 Coefficients of thermal expansion of fibre reinforcements and titanium aluminide matrix alloys determined at various temperatures by thermomechanical analysis [241].
Figure 12 Coefficients of thermal expansion of fibre reinforcements and titanium aluminide matrix alloys determined at various temperatures by thermomechanical analysis [241].
Figure 15 Longitudinal tensile strength of unidirectional SCS-6/ Ti-24Al-11Nb composite versus number of prior 150 to 815°C thermal cycles in air [83]. The even lower inherent “forgiveness” of the y com- positions has precluded the development of meaning-
Figure 15 Longitudinal tensile strength of unidirectional SCS-6/ Ti-24Al-11Nb composite versus number of prior 150 to 815°C thermal cycles in air [83]. The even lower inherent “forgiveness” of the y com- positions has precluded the development of meaning-
Figure 16 Hydrogen activity for selected titanium alloys at 800°C [215, 216].
Figure 16 Hydrogen activity for selected titanium alloys at 800°C [215, 216].
“This involves chemistry and microstructure (and fibre/coating): 10: fully defined, 1: optimum parameters presently unknown. » Terrestrial or non-life-dependent flying application. * Critical life-dependent flying application. 4 Demonstration use rather than utilizing novel characteristics of composite.
“This involves chemistry and microstructure (and fibre/coating): 10: fully defined, 1: optimum parameters presently unknown. » Terrestrial or non-life-dependent flying application. * Critical life-dependent flying application. 4 Demonstration use rather than utilizing novel characteristics of composite.
Figure 18 Comparison amongst titanium-rich alloys, titanium aluminides, superalloys and titanium aluminide composites: (a) yield stress/ density versus temperature, (b) tensile elongation versus temperature, (c) Larson—Miller parameter normalized for rupture life and (d) fracture toughness versus temperature [83]. LMP = T (K) {20 + log[t(h)] }/1000. The present discussion will follow the path taken by Larsen et al. [83] and only consider monolithic and Composite «, material can be fabricated by a num- ber of techniques, although incompatibility problems (CTE differences and chemical reaction) remain with current fibre-matrix material concepts. Assuming
Figure 18 Comparison amongst titanium-rich alloys, titanium aluminides, superalloys and titanium aluminide composites: (a) yield stress/ density versus temperature, (b) tensile elongation versus temperature, (c) Larson—Miller parameter normalized for rupture life and (d) fracture toughness versus temperature [83]. LMP = T (K) {20 + log[t(h)] }/1000. The present discussion will follow the path taken by Larsen et al. [83] and only consider monolithic and Composite «, material can be fabricated by a num- ber of techniques, although incompatibility problems (CTE differences and chemical reaction) remain with current fibre-matrix material concepts. Assuming
Figure 20 (a) High-pressure turbine stator inner support rings for a GE demonstrator engine, run for 25 h and (b) high-pressure compressc casing section, P & W F100 engine, run for 65 h [27].
Figure 20 (a) High-pressure turbine stator inner support rings for a GE demonstrator engine, run for 25 h and (b) high-pressure compressc casing section, P & W F100 engine, run for 65 h [27].
Figure 19 TiAl cast turbocharger rotor [320].
Figure 19 TiAl cast turbocharger rotor [320].
Figure 21 Turbine blade produced from TiAl using a powder metallurgy and hot forging approach [124].
Figure 21 Turbine blade produced from TiAl using a powder metallurgy and hot forging approach [124].
Figure 23 Honeycomb panel constructed of Ti-24Al-11Nb face sheets and Ti-3AI-23V core [329].
Figure 23 Honeycomb panel constructed of Ti-24Al-11Nb face sheets and Ti-3AI-23V core [329].
Figure 22 The weight savings which can be achieved (up to 75%) by replacing a conventional disc and spacer assembly with a titanium aluminide composite ring [281]. oriented so that the large hoop stresses are accom- modated by these filaments. The dramatic weight saving which can be achieved (up to 75%) by replacing
Figure 22 The weight savings which can be achieved (up to 75%) by replacing a conventional disc and spacer assembly with a titanium aluminide composite ring [281]. oriented so that the large hoop stresses are accom- modated by these filaments. The dramatic weight saving which can be achieved (up to 75%) by replacing
Related topics:
EngineeringMaterials ScienceEnvironmental EffectFracture Mechanic
580 California St., Suite 400
San Francisco, CA, 94104
© 2025 Academia. All rights reserved