Monotonic and low cycle fatigue behavior of an O+B2 alloy at high temperatures

Materials Science & Engineering A 599 (2014) 268–278

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Monotonic and low cycle fatigue behavior of an Oþ B2 alloy at high temperatures G. Srinivasulu a, P. Ghosal a, N. Singh c, L. Nazé b, T.K. Nandy a, V. Kumar a, V.V. Kutumbarao a, D. Banerjee d, J.L. Strudel b,n a

Defense Metallurgical Research Laboratory, Kanshanbagh, Hyderabad 500 058, India Centre des Materiaux, CNRS UMR 7633, Mines-ParisTech, BP 87, 91003 Évry Cedex, France c Rm #19, HEPP Building, National Physical Laboratory, New Delhi 110 012, India d Department of Materials Engineering, Indian Institute of Science, Bangalore 560 012, India b

art ic l e i nf o

a b s t r a c t

Article history: Received 25 November 2013 Received in revised form 28 January 2014 Accepted 28 January 2014 Available online 5 February 2014

Low cycle fatigue behavior of an O þB2 alloy was evaluated at 650 1C in ambient atmosphere under fully reversed total axial strain controlled mode. Three different microstructures, namely equiaxed O plus aged B2 (fine O plates in B2 matrix), lenticular O laths plus aged B2 and a pancake composite microstructure comprising equiaxed α2, lenticular O and aged B2, were selected to study the effect of microstructure on low cycle fatigue behavior in this class of alloys. Distinct well-defined trends were observed in the cyclic stress–strain response curves depending on the microstructure. The cyclic stress response was examined in terms of softening or hardening and correlated with microstructural features and dislocation behavior. Fatigue life was analyzed in terms of standard Coffin–Manson and Basquin plots and for all microstructures a prevailing elastic strain regime was identified, with a single slope for microstructures equiaxed and composite and a double slope for lenticular O laths. & 2014 Elsevier B.V. All rights reserved.

Keywords: Intermetallics Plasticity Fatigue High temperature

1. Introduction It has been more than three decades since the initial compositions of Ti3Al base alloys were evaluated for room and high temperature properties [1–3]. A variety of alloying additions were attempted in order to improve room temperature ductility and toughness of these alloys. Additions of Nb showed promise [4] and subsequently several Ti3Al–Nb alloys were developed [5–7] with increasing Nb levels. The metallurgy of these alloys has been reviewed [8–11] with emphasis on tensile and creep properties. The superior properties of alloys rich in Nb have been attributed to the presence of an orthorhombic ordered derivative (designated the O phase) of the hexagonal intermetallic, Ti3Al in these alloys [12,13]. The reasons are twofold: (1) higher solid solution strengthening of the O phase [14,15] and (2) the presence of additional ‘2cþa’ slip [16] that imparts higher elongation to failure and fracture toughness. In 2003, a review of physical, chemical and mechanical properties of this family of alloys [17] concluded that despite their lesser impact in the current literature, orthorhombic titanium aluminides show a well balanced property profile without significant drawbacks. Since then, new efforts have been

n

Corresponding author. E-mail address: [email protected] (J.L. Strudel).

http://dx.doi.org/10.1016/j.msea.2014.01.086 0921-5093 & 2014 Elsevier B.V. All rights reserved.

directed towards property optimization of different O þB2 alloys with the objectives of optimizing strength, toughness and creep. Minor additions of boron (0.1 at%) improve significantly the resistance to HCF crack initiation [18]. Quaternary additions of Zr, Mo, Nb and Si have been attempted with beneficial effects [19–21]. However studies on the low cycle fatigue of this class of alloys, while important for applications, are limited [22–25] and focused more on high cycle fatigue or fatigue crack initiation of intermediate Nb content alloys rather than the more recent high Nb alloys. Therefore, the present study attempts to evaluate monotonic and low cycle fatigue behavior of Ti–20Al–25Nb–1Mo with high Nb content (at%) that exhibits two-phase O þB2 or three-phase O þ α2 þB2 microstructures depending upon the heat treatment.

2. Experimental The alloy (composition shown in Table 1) was melted using double vacuum arc remelting (VAR). This was followed by low strain rate hot die forging at approximately 1150 1C and then by groove rolling at approximately 950 1C, to obtain 14 mm diameter rods. The as-rolled microstructure is shown in Fig. 1. Heat treatments were designed (Table 2, Fig. 2) to obtain three different types of microstructures that were subsequently evaluated for low

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cycle fatigue (LCF) behavior: Er with equiaxed O-phase particles, Hm with O-laths and TNS3 a mixture of both are described in detail in the next section. The monotonic tensile stress–strain curves at 650 1C were carried out in air on an Instron machine, at a constant strain rate of 6  10  4 s  1. The tensile specimens, with a usable gauge length of 25 mm, were equipped with a longitudinal LVDT extensometer. Fatigue tests were performed at 650 1C in air in a servohydraulic machine (100 kN, MTS 880) under fully reversed (Rε ¼  1) total axial strain control mode at a frequency of 0.2 Hz. Mirror-polished cylindrical specimens of 12 mm uniform gauge length and 6 mm diameter were used for the investigation. Typically, three samples per condition were used for tensile testing, while one was used for LCF tests. Data in Fig. 5 is plotted in the form of normalized stress F/S (or true stress) and logarithmic strain in % (or true strain in %) defined as ε ¼100 ln(1 þ e), where F is the applied force, S is the average actual section of the sample S¼ S0(1  e), e is the engineering elongation e¼ Δl/l0 with S0 and l0 respectively the initial cross section area and initial gauge length of the sample. Dislocation substructures were studied by TEM on foils taken from completed fatigue tests, in an attempt to correlate overall dislocation substructures to the fatigue behavior. Two thin foils were prepared from each sample and 6–8 regions were examined before selecting the areas best representative of the microstructures observed, first under a g ¼002 reflection, in order to detect the presence of “2c þa” type dislocations and then usually observed and recorded under one or two different {221}

Table 1 Nominalþ analyzed composition (at%) of the alloy.

Nominal Analyzed

Al

Nb

Mo

Ti

O

N

H

20 18.4

24 25.2

1 0.71

Balance Balance

2050 ppm

550 ppm

1300 ppm

Fig. 1. As rolled microstructure showing globular O phase in a continuous matrix of fine transformed B2.

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pyramidal plane reflections (g ¼{221}). Thin sections of the samples were spark cut at 451 to the stress axis, in order to position them favorably with respect to the highest shear stress and also increase the chances of observing active slip system parallel to the foil plane. The foils were finally thinned down by electrochemically polishing in a twin-jet polishing unit with a 6% sulfuric acid solution in methanol cooled down to  50 1C.

3. Results 3.1. Heat treatment and microstructure Heat treatment schedules selected are presented in Fig. 2 and the resulting microstructures, as observed by SEM, in Fig. 3. The beta transus temperature Tβ was estimated at 1040 720 1C. The variation in beta transus results from composition variations across bands of approximately 100 mm width, as a consequence of microsegregation in the solidification process that was not homogenized during the high temperature ingot processing sequence. (Note that the β phase orders into a B2 structure at a temperature that has not been determined for this alloy.) On cooling further, the alloy passes from a two phase α2 þ β(B2) field to a three phase α2 þB2 þO at a temperature Tα that was not determined and finally to a two phase O þB2 field. The TO þ β transus temperature, below which the O phase is stable, in a twophase equilibrium with the B2 matrix, was estimated at 920 75 1C. Subsolvus heat treatment of 900 1C/24 h was selected in order to preserve the finely distributed equiaxed O phase resulting from the hot rolling process in the α2 þ OþB2 domain. The resulting microstructure designated as Er (Fig. 3a) shows equiaxed globules (about 2 μm in diameter) of O phase (confirmed by XRD analysis) in a continuous matrix of aged B2 (called also B2t), formed during the 650 1C aging treatment. The orthorhombic phase shows a darker contrast in backscattered electron image than the B2 or aged B2. Although the microstructure is similar to that of the asrolled condition (Fig. 1), the equiaxed particles coarsened by a factor of 4 during the 900 1C/24 h heat treatment. The details of aged B2 microstructure that consists of fine platelets or needles of O phase in B2 matrix are seen in Fig. 4a. The presence of multiple variants of O phase, which follow the Burgers orientation relationship with B2, results in a typical diffraction pattern as shown in Fig. 4b. A supersolvus heat treatment above the β transus followed by slow cooling to 900 1C and then the low temperature aging treatment resulted in the microstructure Hm (Fig. 3b). This heat treatment was designed to favor the formation of high aspect ratio coarse O laths with the same volume fraction as the equiaxed O phase in Er in a continuous matrix of large grains of B2t. The heat treatment results in a higher volume fraction of O laths along the prior beta grain boundaries than in the grain interior. Examined by TEM, these laths, in the undeformed material (Fig. 4c), just contained a few isolated dislocations and occasionally a pile-up against an O/B2 interface. All of them invisible under (002) type reflection, these dislocations, generally imaged with (221) reflections, had a ¼[100] or an ¼1/2[110] Burgers vector. The aged B2 structure in this case is similar to that observed in heat treatment

Table 2 Heat treatment details. Heat treatment

Designation

Description

Supertransus treatment

Hm

1080 1C/30 min/FC at 1001/h to 900 1C/24 h/AC to RTþ 650 1C/24 h/AC

Subtransus treatment

Er TNS3

900 1C/24 h/AC þ650 1C/24 h/AC 1000 1C/1 h/FC at 2001/h to 900 1C, then at 151/h to 870 1C/24 h AC to RT þ 650 1C/24 h/AC

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Fig. 2. A schematic description of the three heat treatment schedules.

Fig. 3. (a) Equiaxed Oþ aged B2 microstructure after α–β heat treatment Er, (b) microstructure Hm consisting of relatively coarse and high aspect ratio O phase in a continuous matrix of finely transformed B2, and (c) a composite microstructure TNS3 consisting of equiaxed and lath segments. Globular α2 particles (in dark contrast) are dispersed throughout both areas, with a higher density in the equiaxed portions of the structure (see inset).

Er, in that it comprises several variants of thin, Burgers orientation related O laths in a B2 matrix (Fig. 4a). A heat treatment initiated slightly below Tβ (1000 1C/1 h) was carried out in order to restrict the B2 grain size by the presence of

about 5% volume fraction α2 particles. The intent of the heat treatment was to produce a multimodal microstructure containing an equiaxed component of α2, accompanied by fairly coarse O laths (due to the cooling sequence from the primary heat

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Fig. 4. (a) Details of aged B2 showing fine platelets of O phase; (b) three variants of O phase showing a standard Burgers' orientation relationship with B2; and (c) undeformed O-lath in Hm microstructure.

treatment temperature) in a matrix of aged B2. However the heat treatment generated a composite microstructure shown in Fig. 3c and designated as TNS3. Due to the composition variations described earlier, parts of the material that were enriched in the β stabilizers, Nb and Mo, were above the transus at the primary heat treatment temperature and formed a lath microstructure (similar to but finer than that in Hm) while those regions that were below the transus showed the expected microstructure consisting of globular α2 phase (in dark contrast), coarse O laths and an aged B2 matrix similar to, but slightly coarser than that of Er.

3.2. Monotonic stress–strain behavior The monotonic tensile stress–strain curves, at 650 1C of the three microstructures, are presented in Fig. 5 together with that of the Er microstructure at 550 1C and 25 1C. The Er microstructure with its fine equiaxed O particles embedded in an aged B2 matrix has the highest structural hardening of the three microstructures examined. It has the highest yield strength at all temperatures. At 25 1C, it strain softens just after an initial stress peak and later on strain hardens slightly as strain is increased up to a low rupture elongation of only 3%. At 550 1C, recovery processes are activated and microstructure Er no longer exhibits any initial stress peak but its strain to rupture reaches 14%. At 650 1C, recovery is even more rapid and the material shows strain softening, yet its ductility is preserved and necking appears after 13% elongation. Microstructure Hm is weaker in yield stress but keeps its hardening rate up to 3% strain. Fractographs show prior β grain boundary fractures and numerous cracks at the laths interfaces,

Fig. 5. Monotonic tensile stress–strain curves of the three microstructures at 650 1C.

while the thin ligaments of B2t matrix provide only locally limited ductility. The composite microstructure TNS3 is the weakest in yield stress. It strain hardens at a rate similar to that of Hm initially and exhibits a failure strain of 6.5%, just intermediate between that of Er and that of Hm.

3.3. Cyclic stress response behavior The cyclic stress response curves for the three microstructures are shown in Fig. 6. The degree of softening or hardening is shown

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Fig. 6. Cyclic stress–strain response of the alloy in three different microstructures at 650 1C for different strain amplitudes. (a) Er, showing cyclic softening at higher strains, (b) Hm, showing cyclic hardening at higher strains, and (c) TNS3 showing initial softening, a plateau and hardening in the later half of the cycles.

in Fig. 7 in which a dimensionless ratio ΔΣ n =Δs1 ¼ ððΔsn  Δs1 Þ= ðΔs1 ÞÞ, defined as the difference between the stress amplitude at a cycle n and the stress amplitude of the first cycle, relative to that first cycle amplitude, has been plotted as function of the number of elapsed cycles. Distinct, well-defined trends are observed in the cyclic stress response curves. For the microstructure Er (Fig. 6a) the stress amplitudes are observed to remain stable till fracture, up to an applied total strain amplitude, 7 Δεt/2, of 70.75%. At higher strain amplitudes, considerable cyclic softening is observed. In case of microstructure Hm (Figs. 6b and 7c), stress response remains stable at low strain amplitudes 70.3%. This microstructure initially strain softens at Δεt/2 ¼ 70.5% and finally work hardens after 200 cycles; at higher strain amplitude cyclic hardening is observed from the start. In case of TNS3 (Figs. 6c and 7c), a mixed stress response is recorded. Cyclic softening prevails during the initial 5–10% of lifetime. Table 3 shows the plastic component of the cyclic strain at halflife (Nf/2) and the corresponding stress amplitude; this data is used to examine cyclic stress–strain behavior in Fig. 8. The tensile

yield stress at 650 1C and 0.2% plastic strain for the three microstructures are also plotted for comparison. The microstructure Hm shows very high work hardening at low strain amplitudes but lower work hardening similar to the other two structures at higher strain amplitudes. The fatigue life data in Table 3 for the three microstructures is shown in Fig. 9 as a function of total strain amplitude. The equiaxed structure, Er, shows the best fatigue life at low strain amplitudes, with a crossover to lower fatigue life at higher strain amplitudes. The cyclic behavior in terms of elastic and plastic strain amplitudes is shown in Fig. 10 in the Basquin relationship and Coffin–Manson relationship for the three structures. The Basquin coefficient for Er and TNS3 microstructures is nearly equal (bE  0.10) whereas the microstructure Hm exhibits two-stage sensitivity: one similar to the other two microstructures below Δεt/2 ¼ 70.5% and a higher slope (bE  0.81) at high strain amplitude. The C–M plots are observed to be nearly parallel (cE  0.40) for microstructures Er and TNS3; Hm shows two regimes of behavior one with cE  0.88 at high strain amplitude (Δεt/2 4 70.5%) and one with a small slope (cE  0.15) at small strain amplitude. The variation of inelastic strain with elapsed

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Fig. 7. Variations in degree of hardening and softening with elapsed cycles at 650 1C and different strain amplitudes for the three microstructures: (a) Er; (b) Hm; and (c) TNS3.

Table 3 LCF results for the three microstructures at 650 1C. 7 Δεt/2

Er

Hm

TNS3

Nf

Nf

MPa

cycles

7 Δεp/2 (Nf/2) %

Δsa/2 (Nf/2)

cycles

7 Δεp/2 (Nf/2) %

Δsa/2 (Nf/2)

%

MPa

cycles

7 Δεp/2 (Nf/2) %

MPa

1.0 0.85 0.75 0.65 0.5 0.3

111 424 1026 2543 20,006 –

0.206 0.097 0.053 0.043 0.018 –

916 846 730 696 550 –

– 523 823 – 1198 25,127

– 0.178 0.132 – 0.050 0.029

– 713 633 – 523 321

– 478 – 1462 16,382 –

– 0.183 – 0.102 0.048 –

– 631 – 571 471 –

Nf

cycles is shown for the three microstructures in Fig. 11, for a strain amplitude of 70.85%. TNS3 and Hm microstructures exhibit higher inelastic strain values than the Er microstructure and at a lower stress level than Er, but strain harden. On the contrary, microstructure Er exhibits initial plastic strain amplitude five times smaller than that of the other two microstructures but only twice smaller at the end of the test, due to strain softening.

Δsa/2 (Nf/2)

3.4. Deformation substructure Fig. 12a and Fig. 12c show the deformation substructure of respectively microstructures Er and Hm deformed at a strain amplitude of 70.3%. Planar slip largely dominates as clearly evidenced when a slip plane is observed edge on (Fig. 12c) and pile-ups of dislocations on various slip systems are observed in

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Fig. 8. Cyclic stress–strain behavior for the three microstructures at 650 1C at Nf/2, open symbols for 0.2 % monotonic tensile yield stress. Fig. 11. Plastic strain amplitude as a function of elapsed cycles for the three microstructures at total strain amplitude of 7 0.85%.

amplitude (7 0.3) specimen. However, at higher strain amplitudes, a much higher density of dislocations is seen in between the slip bands (Fig. 12d). Interaction of gliding “a ¼[100] or an ¼1/2 [110]” dislocations (invisible with g¼ 002) with uniformly spaced “2cþ a”slip bands in a single O lath is seen in Fig. 13. The long a and/or a* dislocations gliding in a prismatic plane nearly parallel to the plane of the foil appear to be bowing out between the “2cþ a” slip bands.

4. Interpretation and discussion

Fig. 9. Cyclic life as a function of total strain.

Fig. 10. Combined Basquin and Coffin–Manson plots for different microstructures.

equiaxed O nodules of Er (Fig. 12a). Notice that slip planes are observed inclined in Fig. 12c rather than edge on. Dislocations significantly increase in density with strain amplitude, and subgrain structure form within equiaxed O phase at higher strain amplitudes (Fig. 12b). In the lath microstructure Hm (Fig. 12c), regions in between the slip bands are relatively free of dislocations for the lower strain

We have described the monotonic and fatigue response at 650 1C of an orthorhombic titanium aluminide alloy heat treated to three different microstructures (Fig. 3): a duplex microstructure consisting of fine equiaxed O phase in a B2t matrix (Er), intertwined O laths in a B2t matrix (Hm) and a composite structure (TNS3). The key features of the mechanical response of these microstructures are as follows: (a) Under monotonic loading (Fig. 5), the Er structure has the highest yield strength and shows strain softening after initial hardening at 650 1C, while the Hm structure shows work hardening. The TNS3 structure has the lowest yield strength. The true fracture strain is the highest in Er followed by TNS3, and lowest in Hm. (b) In fatigue, the Er structure shows stable behavior till fracture at an applied total strain amplitude, 7 Δεt/2, of 70.75% and strain softening at higher strain amplitudes (Figs. 6 and 7). The Hm structure (Fig. 7b) shows stable behavior at low strain amplitudes (Δεt/2 ¼ 70.3%), initial softening at 70.5% for the first 200 cycles followed by hardening, and continuous hardening at higher strain amplitudes. The TNS3 structure shows initial strain softening followed by hardening at all strain amplitudes (Fig. 7c); the turning point of the curve takes place after 1000 cycles at 70.5%, after 50 cycles at 70.65% and 5 cycles at 70.85%. In general fracture intervenes for most cases, in these materials before cyclic stability is reached. (c) The Er structure provides the best fatigue resistance at low strain amplitudes but a crossover in behavior with Hm appears at higher strain amplitudes (Fig. 9). We discuss these features in more detail below.

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Fig. 12. Deformation substructures of fatigue tested sample. (a) and (b) correspond to equiaxed microstructures following α–β heat treatment (Er) whereas (c) and (d) show substructure in β treated (Hm) specimens. (a) Pile-ups of dislocations observed sideways in equiaxed O phase (Δεt ¼ 7 0.3%); (b) subgrain formation in equiaxed O phase (Δεt ¼ 7 0.85%); (c) slip bands of “2cþ a” dislocations in O laths (Δεt ¼ 70.3%); and (d) high dislocation density in between the slip bands in O laths (Δεt ¼ 70.85%).

4.1. Monotonic behavior The monotonic behavior of the Er and Hm microstructures is determined initially by the relative strengths and work hardening capacities of the primary equiaxed O and lath O constituent respectively because these are expected to be softer than aged B2 with its associated nanometric O phase dispersion (Fig. 4). The primary O lath structure in Hm has a significantly larger length scale (an order of magnitude) in at least one direction as compared to equiaxed primary O phase in the Er structure (Fig. 3). It is therefore expected to be softer than the equiaxed O through a size effect, since for some specific slip systems and dislocation line directions, the Orowan bowing out can extend to the length of the O laths (Fig. 13). Thus the yield strength of the Hm structure is lower than that of the Er structure and it strain hardens to higher strain amplitudes than Er and to similar peak strength levels, i.e. 830 MPa. In HCP crystals, slip systems operate rather independently and the formation of attractive junctions and cell walls requires higher strain levels than in crystals with higher symmetry, such as B2, fcc or bcc crystals. The role of dislocation pile-ups

forming in prismatic, basal and pyramidal planes in titanium and its alloys becomes predominant in the behavior of these alloys. This planar slip activity is also observed in this ordered phase, especially at lower strain amplitude (Figs. 12a, c, d and 13). Whenever either, within the equiaxed O of Er or within the laths O of Hm, dislocation pile-ups generate local stress concentration levels capable of inducing plasticity in the surrounding B2 transformed matrix, then slip transfer can occur into that constituent of the structure. As shown in Fig. 5 the tensile curve of Er at room temperature suggests that in Er, the fine, equiaxed primary O phase work hardens rapidly from an initially high yield stress to the level where slip transfer to the aged B2 is taking place, just after the stress peak at RT, for instance. At higher temperature, work softening of this highly strengthened constituent ensues, probably by the sudden shearing and disordering of the many variants of ultrafine tertiary O laths (Fig. 4a); then the subsequent work hardening resulting from the multiplication and interactions of slip systems in the aged B2 is rather low but effective at RT. This behavior of the B2t matrix is similar to that observed in beta alloys of titanium containing fine dispersions of α phase [26]. Indeed, at

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corresponding to the kinetic balance between strain hardening at the selected strain rate and thermally activated recovery of these microstructures at 650 1C. In this case fracture intervenes before significant plastic deformation can take place, since fracture is initiated by early intergranular cracking of the large prior β grains generated during the supersolvus treatment at 1080 1C/ 0.5 h [27]. The TNS3 structure appears to combine features of both the Hm and Er structures in that it has low yield strength but significant extent of plastic deformation with low work hardening. Notice its UTS (730 MPa) lower than that of the other two microstructures, a weakness that could result from a B2t matrix leaner in tertiary O laths, as a result of the last part of the first heat treatment carried out at 870 1C, rather than at 900 1C (Fig. 2). 4.2. Cyclic behavior

Fig. 13. Microstructure Hm: interaction of a/a* dislocations with 2c þ a dislocations in an edge-on pyramidal plane of an O lath (Δεt ¼ 7 0.85%). The long wavy dislocation segments alone are invisible with g ¼ 002, indicating that they have a¼ [100] or a* ¼1/2[110] Burgers vectors, gliding on a prismatic plane (shaded on oriented prism), nearly parallel to the plane of the foil.

higher temperatures plasticity is not only mechanically activated but also thermally activated and recovery mechanisms become operative. Consequently at 550 1C, the Er microstructure tested at a strain rate of 6  10  4 s  1 (Fig. 5) exhibits a delicate and exact balance between the hardening produced by straining and the softening provided by recovery processes. At 650 1C, under the same strain rate, Er actually work softens, suggesting that now, recovery processes are prevalent over strain hardening mechanisms within that structure. In contrast, the Hm structure, at 650 1C, work hardens more gradually to higher strain levels from an initially lower yield stress to about the same peak stress as Er, i.e. 830 MPa, a level

We may now discuss cyclic behavior against the background of monotonic plasticity in these structures. As is well known, materials that have high UTS in relation to yield strength ( Z1.2) tend to cyclically harden while the reverse is true for materials with a low ratio of UTS to yield stress [24,28]. This general trend is reflected in the behavior of this aluminide alloy as well. The Er structure with its low UTS/sy ratio (1.15) cyclically softens while the reverse is true for the Hm structure with a ratio of 1.26. Looking more finely at the data presented in Figs. 6 and 7, we notice that when the plastic strain is extremely low (2–3  10  4), the Hm and Er microstructures are remarkably stable cyclically. It is a range of plastic strain amplitude rarely explored previously, where plastic strain is confined to a limited number of planar slip systems fed by activated dislocation sources generating shear loops that pileup against interfaces. These planar pileups have been observed by Brandes et al. [30] in a Ti–6Al strained at RT, who report the absence of any measurable backflow, at this temperature, upon unloading due to age hardening phenomena. On the other hand, Jousset [31] repeatedly observed pileups stored against laths boundaries in Ti 6242 and interpreted the reverse straining recorded during cyclic loading at 300 1C and above, as produced by the reversal movement and cancelation of these dislocations bowing out and multiplying events that generate pileups. DSA and solute segregation effects can no longer immobilize pileups at these temperatures where the mobility of solute elements has increased significantly. Thus plastic straining at low strain amplitude appears to be completely reversible, when confined to the O phase where both crosslip events and attractive junctions between dislocations are absent; as a result, cycling plastically at Δεp/2 ¼ 70.03% would not produce any hardening or softening of the material. At higher strain amplitude, on the other hand, when plastic strain extends into the B2 transformed matrix under the enhanced stress field produced by the pileups, strain softening is observed. Cyclic softening is also observed in the disordered equivalent of the O phase alloys that is in titanium alloys with bimodal structures containing primary alpha and aged beta [29,32]. In this case softening has been shown to occur in stages, a first rapid softening due to the shearing of fine, ordered Ti3Al particles in the α phase by planar slip bands, and a second stage due to destruction of the α⧸β boundaries in the aged β matrix. The second of these mechanisms is clearly similar in both ordered and disordered alloys. A comparison of Fig. 5 with Fig. 6 suggests that cyclic softening in Er initiates at strain amplitudes at which work softening begins in the monotonic curves, that is at stress levels at which slip is transferred to aged B2. The softening process must therefore be related to destruction of the fine aged B2 microstructure by a combination of repeated shearing by dislocations and thermal processes, leading to reduced strength of obstacles for

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dislocation motion. Similarly, when Hm is tested at Δεt/2 ¼ 7 0.5% and plastic straining reaches critical level of 0.05% (Fig. 10), cyclic softening can be observed and persist during the first 200 cycles, although it was not observed during tensile testing. This peculiar point can be rationalized by observing that low amplitude cyclic testing can accumulate large amounts of strain in small volumes by preserving strain localization, when tensile testing will necessarily combine large amounts of strain with generalized plasticity due to strain hardening. Finally, at high strain amplitude in microstructure Hm, multiple slip systems are shearing all phases, interact with each other thus generating substructures (Fig. 12b) responsible for strain hardening. This feature of deformation behavior seems to be associated with a change in cyclic work hardening behavior as shown in Fig. 8. As shown in Figs. 9 and 10 this transition from O confined plasticity and reversible plasticity operating at very low strain amplitude to generalized plasticity and work hardening operating at high strain amplitudes is marked by a drastic change in slope, with that of the high strain regime rather similar to the value of c E  0.82 reported for a lath microstructure tested at 650 1C [24]. Yet, during this transition from localized to generalized plasticity the materials experience a brief strain softening stage when planar slip systems extend into the aged B2 matrix and penetrate, shear and disperse the many thin variants of tertiary O phase. This short intermediate stage is present in Hm at the turning point (Fig. 7b) where Δεt/2 ¼ 7 0.5% and Δεp/2 reach the critical value of 4  10  4 (Fig. 8): strain softening prevails during the first 200 cycles when local straining is spreading out into the B2 matrix and finally, generalized plasticity takes over during the last 20,000 cycles. Similarly, microstructure TNS3 (Fig. 7c), which contains a wide spectrum of O phase morphologies and sizes, ranging from micron size equiaxed nodules to 15–20 mm long needles, experiences this intermediate stage of softening before consolidating for the rest of the fatigue lifetime of the material. In the narrow range of total strain amplitude explored (0.5–0.85) the transition from softening to hardening happens sooner and with a lesser stress amplitude as strain is increased (quantitative values are gathered in Table 4). One might infer that TNS3 would start hardening right from the first cycles for strain amplitudes of 1% as observed with Hm microstructure in its high strain regime since all possible O phase (lath or nodule) present in the material up to the smallest equiaxed nodule will contribute immediately to general plasticity. 4.3. Fatigue life Fig. 9 shows that the Er structure has better fatigue resistance at low strain amplitudes. The TNS3 behavior in terms of fatigue life is equivalent to Er. In order to understand this behavior, the total strain has been decomposed into its elastic and plastic components as a function of cycles to failure in Fig. 10. At low strain amplitudes, fatigue is governed by a macroscopic elastic strain response in all three structures, a usual feature for this family of highly strengthened alloy as already reported for a 24Al–15Nb– 1Mo alloy (22). In this regime the Hm structure has a fatigue strength exponent that is almost twice that of Er and TNS3. This reflects the higher fatigue work hardening of the Hm structure (Fig. 8) (the strength exponent is given approximately by b¼  n0 / Table 4 Softening to hardening transition parameters for TNS3 microstructure.

Δεt/2 Δεp/2 Nb of cycles ΔΣ n =Δs1 Relevant size O phase

0.5% 0.05% 1 000 0.07 Above 20 mm (  Hm)

0.65% 0.12% 50 0.045

0.85% 0.2% 5 0.03 Down to 1 mm (  Er)

277

Fig. 14. Strain controlled fatigue life of the current alloy and microstructural modifications in comparison with a conventional titanium alloy, IMI 834 and other titanium aluminides. Data sources are indicated in the figure legend.

(1 þ5n0 ) [28]). The higher fatigue strength exponent of the Hm structure leads to its lower fatigue life in the lower strain amplitude regime. At higher strain amplitudes the Hm structure shows a transition in behavior such that the trends suggest that its fatigue life resistance becomes superior to Er (Fig. 9). This transition is also reflected in a sharp change in slope in the Coffin– Manson plot (Fig. 10) and in cyclic work hardening behavior (Fig. 8). We note that in this range of strain amplitudes, Hm shows cyclic hardening rather than cyclic softening in contrast to Er (Fig. 6), and the sharp change in trend in fatigue life with increasing strain amplitude of the Hm structure may reflect this behavior. We note in this context that extrapolation of the Coffin– Manson slope to higher strain levels in Er indicates that transition from elastic strain controlled to plastic strain controlled cyclic behavior will occur at about 10  2 plastic strain amplitude (Fig. 10), at which the life will be less than one-quarter cycle (tensile portion of the hysteresis loop). Therefore LCF testing at much higher strain amplitude (4 1.0%) is not possible. Finally we compare the fatigue response of this orthorhombic titanium alloy with that of similar alloys with lower niobium [22,24] which contain larger volume fractions of the α2 phase (Fig. 14) as well as with a conventional Ti alloy IMI-834 [29]. The strain controlled fatigue life of the orthorhombic alloy of this study, in its different microstructural forms, covers the range of fatigue lives obtained for the alloys with lower niobium. However for the comparable microstructure (Er) the orthorhombic titanium alloy offers better fatigue life, especially at lower strain amplitude, probably because cleavage crack initiation in the fine equiaxed particles of the O phase is likely to occur at higher local strain levels than in equiaxed particles of the α2 phase [32]. The data for the conventional alloy IMI-834 was obtained at 600 1C.

5. Conclusion 1. The monotonic response and the strain controlled fatigue response of an orthorhombic titanium aluminide alloy were investigated as a function of microstructure at 650 1C. 2. A bimodal structure with equiaxed O primary phase has higher yield strength in comparison to the bimodal structure with lath O primary phase. The elongation to failure of the equiaxed microstructure is significantly higher than that of the lath structure.

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3. The equiaxed structure shows work softening at high strain amplitudes in strain controlled fatigue, while the lath structure shows work hardening. Fracture intervenes before either structure reaches a stable hardening response. 4. Reasonable fatigue lives are realized in all structures only when the elastic strain component of the total strain dominates the strain controlled fatigue response. The equiaxed structure shows superior fatigue life at low strain amplitudes, possibly due to its larger elastic response to the mechanical solicitation.

Acknowledgments The authors acknowledge the contribution of the Indo-French Center for the Promotion of Advanced Research (CEFIPRA) which offered a 2 year post-doctoral fellowship to one of us (Dr. N. Singh). We acknowledge the financial support of DRDO through the Defense Metallurgical Research Center in Hyderabad (India), of SAFRAN Group through Snecma Moteurs and Mines-ParisTech in Paris (France).

[9] [10] [11]

[12] [13] [14] [15] [16] [17]

[18] [19]

[20] [21]

References [1] S.M.L. Sastry, H.A. Lipsitt, Metall. Trans. A 8 (1977), p. 1543. [2] P.L. Martin, H.A. Lipsitt, N.T. Nuhfer, J.C. Williams, in: H. Kimura, O. Izumi (Eds.), Titanium'80, Science and Technology, Metallurgical Society-A.I.M.E., PA, USA, 1980, p. 1245. [3] H.A. Lipsitt, High temperature ordered intermetallic alloys, MRS 39 (1985), 352. [4] J.B. McAndrew, C.R. Simcoe, Investigation of the Ti–Al–Cb System as Source of Alloys for Use at 1220–1800 1F, Armour Research Foundation of Illinois, Institute of Technology, WADD-TR-60-99, Wright Patterson Airforce Base, 1960. [5] M.J. Blackburn, M.P. Smith, Titanium alloys of the Ti3Al type, US Patent 4.292.077, 1981. [6] M.J. Blackburn, M.P. Smith, Titanium aluminum alloys containing niobium, vanadium and molybdenum, US Patent 4.716.020, 1987. [7] R.G. Rowe, US Patent No. 032, 357, 1991. [8] D. Banerjee, A.K. Gogia, T.K. Nandy, K. Muraleedharan, R.S. Mishra, in: R. Darolia, J.J. Lewandowski, C.T. Liu, P.L. Martin, D.B. Miracle, M.V. Nathal,

[22] [23] [24] [25]

[26] [27] [28] [29] [30] [31] [32]

Proceedings of the 1st International Symposium on Structural Intermetallics, TMS, 1993, pp. 19. D. Banerjee, in: J.H. Westbrook, R.L. Fleischer (Eds.), Intermetallic Compounds – Principles and Practice, 2, John Wiley & Sons Ltd., England, 1995, p. 91. A.K. Gogia, T.K. Nandy, D. Banerjee, D.T. Carisey, J.L. Strudel, J.M. Franchet, Intermetallics 6 (1998) 741. D. Banerjee, R.G. Baligidad, A.K. Gogia, J.L. Strudel, in: K.J. Hemker, D.M. Dimiduk, H. Clemens,R. Darolia, H. Inui, J.M. Larsen, V.K. Sikka, M. Thomas, J.D. Whittenberger (Eds.), TMS (The Minerals, Metals and Materials Society), 2001, p. 43. D. Banerjee, A.K. Gogia, T.K. Nandy, V.A. Joshi, Acta Metall. 36 (1988), p. 871. D. Banerjee, Prog. Mater. Sci. 42 (1) (1997) 135. T.K. Nandy, PhD thesis, Banaras Hindu University and Defence Metallurgical Research Laboratory, India, 1999. T.K. Nandy, D. Banerjee, Intermetallics 8 (2000) 1269. D. Banerjee, Philos. Mag. A 72 (1995) 1559. J. Kumpfert, C. Leyens, in: C. Leyens, M. Peters (Eds.), Titanium and Titanium Alloys. Fundamentals and Applications, Wiley-VCH, Weinheim, Germany, 2003, pp. 59–88. M. Hagiwara, T. Kitaura, Y. Ono, T. Yuri, T. Ogata, S. Emura, Mater. Trans. JIM 53 (6) (2012) 1138–1147. L. Germann, D. Banerjee, J.Y. Guedou, J.-L. Strudel, in: G. Luetjering, J. Albrecht (Eds.), Ti-2003, Science and Technology, Wiley, Weinheim, Germany, 2004, pp. 2137–2144. L. Germann, J.Y. Guedou, D. Banerjee, J.L. Strudel, Intermetallics 13 (2005) 920–924. T. Carisey, D. Banerjee, J.M. Franchet, A.K. Gogia, A. Lasalmonie, T.K. Nandy, J.L. Strudel, Titanium based intermetallic alloys, US Patent No. 6.132.526, 2000. P.R. Bhowal, W.A. Konkel, H.F. Merrick, Mater. Sci. Eng. A 92/193 (1995) 891–898. P.N. Singh, B.K. Singh, C. Ramachandra, Vakil Singh, Scr. Mater. 34 (11) (1996) 1791–1796. J. Cao, F. Bai, Z.h. Li, Mater. Sci. Eng. A424 (2006) 47–52. S. Lutjering, P.R. Smith, D. Eylon, EUROMAT 1999, Biannual Meeting of Federation of European Materials Societies (FEMS), Munich, vol. 10, 2000, pp. 283–288. Bhattacharjee, PhD thesis, IIT Kanpur, India, 2006. G. Srinivasulu, N. Singh, L. Naze, T.K. Nandy, V. Kumar, V.V. Kutumba Rao, D. Banerjee, J.L. Strudel, CEFIPRA Project 1908D (Final report), 2005. R.W. Hertzberg, Deformation and Fracture of Engineering Materials, New York, John Wiley and Sons (1996) 556. N. Singh, GouthamaVakil Singh, Int. J. Fatigue 29 (2007) 843–851. M.A. Brandes, M. Mills, MSE-A387-389 (2004) 570–575. H. Jousset, PhD thesis at Centre des Materiaux, Mines-ParisTech, December 2008. S. Hardt, H.J. Maier, H.J. Christ, Int. J. Fatigue 21 (1999) 779–789.