- Email: [email protected]
Property optimization through microstructural control in titanium and aluminum alloys
Materials Science and Engineering A263 (1999) 117 – 126
Property optimization through microstructural control in titanium and aluminum alloys G. Lu¨tjering * Technical Uni6ersity Hamburg-Harburg, 21071 Hamburg, Germany
Abstract In this paper the microstructure/property relationships of Ti-alloys (a + b and b alloys) and high strength, age-hardened Al-alloys are discussed. Common features in these microstructure/property relationships can be found for a+ b Ti-alloys and Al-alloys in the underaged condition as well as for b Ti-alloys and Al-alloys in the overaged condition. Basic differences are the presence of dispersoids and large Fe and Si containing inclusions in Al-alloys and the alloying element partitioning effect in bi-modal structures of a +b Ti-alloys. © 1999 Elsevier Science S.A. All rights reserved. Keywords: Al-alloys; Fe; Si; Ti-alloys; Microstructure; Mechanical properties
1. Introduction
2. Microstructure of a+ b Ti-alloys
Titanium alloys (a + b and b alloys) and aluminum alloys are widely used in aerospace applications because of their low density and high strength values. Typical values for some important mechanical properties are shown in Table 1 for Ti – 6Al – 4V (a + b Ti-alloy), b-CEZ (b Ti-alloy) and the high strength, age-hardened Al-alloy 7475. It can be seen that the a + b Ti-alloy Ti –6Al–4V shows the best resistance to crack propagation whereas the b Ti-alloy b-CEZ exhibits the highest tensile and fatigue strength. Comparing these values to the Al-alloy 7475, the difference in density (4.5 g cm − 3 for Ti-alloys and 2.7 g cm − 3 for Al-alloys, a factor of about 1.7) should be taken into account. For all three types of alloys mechanical properties are optimized to a large extent through microstructural control by appropriate processing and heat treatment. The present paper tries to outline common features in such microstructural control and in the microstructure/property relationships as well as basic differences. Common guide-lines for these microstructure/property relationships might be useful for further improvements in mechanical properties for both Ti-alloys as well as for Al-alloys.
The main microstructural features of fully lamellar structures in a+ b Ti-alloys (Fig. 1) are the relatively large size of the equiaxed b grains (about 600 mm), the presence of continuous a-layers at b grain boundaries, the a-colony size and the size of the individual a-lamellae. Since the recrystallization kinetics in the b phase field are so fast, only the cooling rate from the b phase field has a major influence on microstructure in the whole processing route. With increasing cooling rate the size of the a-colonies as well as the size of the individual a-lamellae are reduced, at extreme fast cooling rates down to the dimensions of a single martensitic plate. For nearly all a+ b Ti-alloys the a phase is age-hardened by the precipitation of very small (nanometer sized), coherent Ti3Al particles (a2) during the final low temperature heat treatment, such as 24 h 500°C (Ti– 6Al–4V), 8 h 590°C (Ti-6242), 2 h 700°C (IMI 834). The volume fraction of these a2 particles depends also on the oxygen content of the alloy, because oxygen is known to stabilize the a2 phase. In contrast to the fully lamellar structures (Fig. 1), a so-called bi-modal (or duplex) type of microstructure (Fig. 2) can be obtained in a+ b Ti-alloys by thermomechanical processing (Fig. 3). In this case the lamellar microstructure is deformed in the a+ b phase field and then recrystallized in the a+b phase field leading to a mixture of equiaxed a and
* Tel.: +49-40-7718-3036; fax: +49-40-7718-2325. E-mail address: [email protected] (G. Lu¨tjering)
0921-5093/99/$ - see front matter © 1999 Elsevier Science S.A. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 9 8 ) 0 1 1 6 9 - 1
G. Lu¨tjering / Materials Science and Engineering A263 (1999) 117–126
118
Table 1 Typical values for mechanical properties at room temperature Alloy
Ti–6Al–4V (bi-modal) b-CEZ (bi-modal) Al-7475 (pancakes)
E (GPa)
sy (MPa)
T.E. (%)
HCF R =−1 (MPa)
DK at 10-9 (m cy.−1)
KIC (MPa m−1/2)
Microcracks R=−1 (MPa m−1/2)
Macrocracks R =0.1 (MPa m−1/2)
115
950
22
500
3
8
60
120
1200
13
625
2
4.5
37
70
480
14
200
1
3.5
45
b grains [1]. Upon cooling from the bi-modal recrystallization temperature the b grains are transformed again to a lamellar structure. The main microstructural difference to fully lamellar structures is the small b grain size (about 20–50 mm) determined by the distance of the primary a grains (ap). This small b grain size of bimodal structures limits the maximum a-colony size and the maximum length of a-lamellae as well as the effective length of the GB a-layers. The most important parameter of the processing route is the correct choice of the recrystallization or bi-modal annealing temperature because it determines the volume fraction of ap. The aging response of bi-modal microstructures is somewhat different from fully lamellar microstructures because of the alloying element partitioning effect taking place during the bi-modal recrystallization treatment and leading to an enrichment in aluminum and oxygen in the ap phase. As a consequence, the lamellar grains in the bi-modal structure are softer (less a2 particles) as compared to the fully lamellar structure.
temperature. Instead, the a phase is precipitated in form of incoherent a-plates within the b grains upon annealing or aging in the a+ b phase field. The main difference between the various microstructures in b Ti-alloys is the size and shape of the b grains (Figs. 4–6). The size of the equiaxed b grains, resulting from recrystallization in the b phase field, is somewhat smaller (about 400 mm) than in a+ b Ti-alloys because of the 150–200°C lower b transus. Upon cooling from the b phase field, extensive GB a-layers are formed (Fig. 4). The heat treatment in the a+ b phase field consists usually of two steps. First, the annealing is performed at relatively high tremperatures leading to
3. Microstructure of b Ti-alloys The characteristic feature of b Ti-alloys is that the content of b stabilizing elements is high enough to suppress martensitic transformation as well as the formation of a a-colony structure upon cooling to room
Fig. 2. a+ b Ti-Alloys, bi-modal structure.
Fig. 1. a+ b Ti-Alloys, lamellar structure.
Fig. 3. a +b Ti-Alloys, processing route for bi-modal structures (schematically).
G. Lu¨tjering / Materials Science and Engineering A263 (1999) 117–126
Fig. 4. b Ti-Alloys, lamellar structure.
119
Fig. 8. b Ti-Alloys, through-transus processing route for necklace structures (schematically).
Fig. 5. b-Ti-Alloys, bi-modal structure.
the formation of large a-plates visible in the LM-picture of Fig. 4. The second step is an aging treatment at lower temperatures (around 600°C) precipitating fine aplatelets shown in the TEM-picture of Fig. 7. The
Fig. 6. b Ti-Alloys, necklace structure.
Fig. 7. b Ti-Alloys, secondary a-platelets.
Fig. 9. Al-Alloys, pancake grain structure.
resulting microstructure of this coarse-grained b recrystallized condition (Fig. 4) is commonly called ‘lamellar’. To avoid the negative influence of the extensive GB a-layers on mechanical properties, two different approaches have been taken for b Ti-alloys [2]. One approach is to obtain a bi-modal type microstructure (Fig. 5) similar to the bi-modal microstructure in a+ b Ti-alloys. The other approach is to create a so-called necklace type microstructure (Fig. 6). The thermomechanical processing route to obtain bi-modal microstructures is similar to the route for a+ b Ti-alloys shown in Fig. 3, except that after the bi-modal recrystallization treatment the above described two-stage heat treatment is used for b Ti-alloys to form large a-plates and small a-platelets. To create a necklace type microstructure (Fig. 6) a so-called ‘through-transus’ deformation process has to be applied (Fig. 8). After homogenization the material is not cooled to room temperature but directly deformed starting in the b phase field and continued into the a+b phase field until a forms at the deformed b grain boundaries as round particles stabilizing the elongated or pancake shaped b grain structure (Fig. 6). It is important that the deformation process is finished before a starts to precipitate also within the interior of the deformed b grains as round particles. Upon cooling, the b grains recover and the final two-step heat treatment forms the desired mixture of large a-plates and small a-platelets (Fig. 7) also in this necklace type of microstructure.
120
G. Lu¨tjering / Materials Science and Engineering A263 (1999) 117–126
4. Microstructure of high strength Al-alloys High strength Al-alloys are age-hardened alloys based on the Al– Cu system (2XXX series) or on the Al–Zn system (7XXX series). Particles of the equilibrium phase, for example Al2Cu (u), are incoherent with the fcc a phase, but upon aging at temperatures below about 200°C small metastable precipitates (size 2 –20 nm) are formed (first coherent u¦ and then semicoherent u%) leading to the characteristic age-hardening curve in which the yield stress is passing through a maximum with aging time. In the underaged condition (before the maximum), coherent u¦ particles are present which can be sheared by the moving dislocations leading to the formation of intense slip bands. In the overaged condition (after the maximum), semi-coherent u% particles are present which cannot be sheared that easily and are mostly by-passed by the moving dislocations (Orowan–mechanism) creating dislocation loops around the particles and a more homogeneous slip distribution. In addition, two other kinds of particles are present in commercial Al-alloys: small incoherent dispersoids (size about 0.2 mm) containing alloying elements such as Mn, Cr, Ti forming intermetallic compounds (Al6Mn, Al3Cr, Al3Ti) and large inclusions (size about 5–10 mm) containing the impurity elements Fe and Si. Both, the dispersoids and the large inclusions are stable up to the melting point. The resulting
Fig. 12. Al-Alloys, dispersoids and subgrains.
Fig. 13. Al-Alloys, processing route for pancake structure (schematically).
microstructural features of high strength Al-alloys are illustrated in Figs. 9–12. As a consequence of the deliberately added, homogeneously distributed dispersoids (Fig. 12) inhibiting recrystallization, high strength Al-alloys exhibit usually an unrecrystallized pancakeshaped grain structure (Fig. 9). The processing route (Fig. 13) consists of homogenization in the single a phase field (at around 500°C) followed by deformation in the temperature region of about 300–400°C (two phase region). During the subsequent solution treatment in the a phase field all precipitated equilibrium phase is re-solutioned but the material is not recrystallizing, only recovery into small subgrains takes place (Fig. 12). The last step in the processing route is the aging treatment to precipitate the coherent or semi-co-
Fig. 10. Al-Alloys, recrystallized grain structure.
Fig. 11. Al-Alloys, PFZ along grain boundaries.
Fig. 14. Al-Alloys, processing route for recrystallized structure (schematically).
G. Lu¨tjering / Materials Science and Engineering A263 (1999) 117–126
121
Table 2 Influence of microstructural parameters on mechanical properties of a+b Ti-alloys and underaged Al-alloys a+b Titanium alloys
sy
Aging (a2), oxygen Bi-modal structure GB a-layers Small a-colonies, alamellae
oF
HCF
Micro-cracks da/dn
Macrocracks
DKth R = 0.7
KIC
DKth R = 0.1
Creep strength 0.2%
Underaged aluminum alloys
+ + O +
− + − +
+ − + − +
− + − +
− − O −
− − − −
+ − O −
+ − O O −
Aging (coh. prec.) Pancakes, small GS PFZ
O
+
+
−
−
−
−
+
O
−
−
O
O
−
O
O
Dispersoids (Mn, Cr, Ti) Inclusions (Fe,Si)
Table 3 Influence of microstructural parameters on mechanical properties of b Ti-alloys and overaged Al-alloys b Titanium alloys
Overaging (asec) GB a-layers Bi-modal structure Necklace structure
sy
oF
HCF
Micro-cracks da/dn
Macrocracks
Creep strength 0.2% Overaged aluminum alloys
DKth R = 0.7
KIC
DKth R = 0.1
− O O
+ − +
− − +
+ − +
+ O −
+ − −
+ O −
− O O
Overaging (semi-coh. prec.) PFZ Small GS
O
+
+
+
−
+
−
O
Pancakes
O O
O −
O −
O O
− O
O −
− O
+ O
Dispersoids (Mn, Cr, Ti) Inclusions (Fe,Si)
herent particles. Unfortunately, particles of the incoherent equilibrium phase, for example Al2Cu (u), are formed during the aging treatment at the grain boundaries creating soft, precipitate-free-zones (PFZ) along the grain boundaries (Fig. 11). The formation of the PFZ is starting already in the underaged stage before the maximum in yield stress is reached but the negative effects of the PFZ are much more pronounced for the overaged condition. In order to obtain recrystallized microstructures with small, nearly equiaxed a grains (Fig. 10), a fairly complicated processing route has to be applied (Fig. 14) [3]. After homogenization the material is annealed in the two phase region to precipitate the equilibrium phase as very coarse particles. The subsequent deformation process is performed at relatively low temperatures followed by recrystallization in the two phase region at temperatures close to the phase boundary. Afterwards a short solution treatment is performed to re-solution the equiblibrium phase followed by the final aging treatment. It should be pointed out that this complicated processing route is not very practical, another way is to reduce the content of the dispersoid-forming elements and to deform the material in the normal processing route (Fig. 13) at as
low as possible temperatures. The resulting grain structure is not as uniform as shown in Fig. 10, the grain size is in most cases fairly inhomogeneous due to the reduced volume fraction of dispersoids.
5. Correlation between microstructure and mechanical properties In this correlation between microstructure and mechanical properties the a+ b Ti-alloys can be compared to Al-alloys in the underaged condition (Table 2), because both materials contain coherent, shearable precipitates, whereas the b Ti-alloys should be compared to Al-alloys in the overaged condition (Table 3), because both materials contain non-shearable precipitates (incoherent a-platelets in case of b Ti-alloys and semicoherent precipitates in case of overaged Al-alloys). In Table 2 and Table 3 the relevant microstructural parameters are listed together with important mechanical properties and the experimentally observed influence is qualitatively indicated by + (positive), − (negative), or O (no influence). Within the scope of this paper, it is not intended to discuss all the individual relationships
122
G. Lu¨tjering / Materials Science and Engineering A263 (1999) 117–126
Table 4 Influence of aging on fracture toughness Ti-alloy Ti–8.6Al
Fracture toughness (MPa m−1/2)
10 h 500°C 260 h 500°C
75 70
[4 –7] in detail. In the following only a rough overview will be given with some examples.
5.1. Comparison of a + b Ti-alloys and underaged Al-alloys Increasing the aging degree in underaged Al-alloys and in a +b Ti-alloys as well as the oxygen content in the latter case results in the same qualitative influence on all mechanical properties listed in Table 2. Increasing the degree of age-hardening by coherent, shearable
Al-alloy X-7075 55 40
24 h 100°C 20 h 160°C
particles will increase the intensity of slip band formation reducing the ductility. The negative influence on fatigue crack nucleation is usually overcompensated by the increase in yield stress. The fatigue crack propagation of small, naturally initiated surface cracks (microcracks) which propagate in this stage in single slip bands (flat crack front) is increased with the increasing intensity of the slip bands. Large macrocracks (CT type of
Fig. 17. Ti-Alloy IMI 834, microcrack propagation (R= − 1). Fig. 15. Ti-Alloy IMI 834, HCF at RT, R = −1, fully lamellar and bi-modal microstructures.
Fig. 16. Al-Alloy X-7075, underaged, HCF at RT, R = −1, influence of grain size.
Fig. 18. Al-Alloy X-7075, underaged, microcrack propagation (R = −1).
G. Lu¨tjering / Materials Science and Engineering A263 (1999) 117–126
Fig. 19. Ti – 6Al – 4V, effect of cooling rate on tensile elongation.
123
erties as changing from fully lamellar structures to bi-modal structures for a+ b Ti-alloys, except for two properties, HCF and creep strength, (Table 2). In these two cases, Table 2 shows in the left column the dependence for bi-modal structures of a+ b Ti-alloys and in the right column the dependence for Al-alloys. This different behavior is demonstrated for the HCF strength by comparing Fig. 15 and Fig. 16 and can be attributed to the alloying element partitioning effect in bi-modal structures of a+ b Ti-alloys leading to crack nucleation in the lamellar regions of the bi-modal structure at lower stress amplitudes than in fully lamellar structures. This alloy element partitioning effect in bimodal structures of a+ b Ti-alloys can also explain the difference between Al-alloys and a+ b Ti-alloys in case of creep strength. It should be pointed out that the negative influence of the alloy element partitioning effect in a+b Ti-alloys can be eliminated by introduc-
Fig. 20. Al-7475, fatigue crack nucleation at large inclusions.
specimens), in the absence of crack closure (high Rvalue, KIC), exhibit lower DKth and KIC values with increasing aging because the crack front profile is dominated by the grain size and is not significantly changed by the intensity of slip bands. The decrease in fracture toughness with increasing degree of age-hardening is illustrated in Table 4 for Ti- and Al-alloys. The fatigue crack propagation threshold value of macrocracks at low R-values is usually increased with increasing aging (Table 2), because the crack closure contribution, increasing with increasing shear displacement at the crack tip (mode II), usually overcompensates the easier fatigue crack propagation with increasing aging (see microcracks). The creep resistance listed in Table 2 (opl. + 0.2%, primary creep) is valid for the dislocationcreep regime (up to 200°C for Al-alloys and 600°C for Ti-alloys). This creep resistance is increasing with agehardening degree (increasing yield stress). Reducing the a grain size or changing to pancake-shaped grains for Al-alloys show the same tendency on mechanical prop-
Fig. 21. Al-Alloys, influence of Fe and Si content on HCF (R = −1).
Fig. 22. b-CEZ, effect of size of a-plates on da/dN– DK curves.
G. Lu¨tjering / Materials Science and Engineering A263 (1999) 117–126
124
Fig. 23. Al-Alloy X-7075, effect of overaging on da/dN– DK curves.
ing in the processing route (Fig. 3) an intermediate heat treatment in the temperature region of 800 – 850°C between the bi-modal recrystallization treatment and the final aging treatment. The alloy element partitioning effect does not play a significant role for all other mechanical properties listed in Table 2. This is illustrated by the example of microcrack propagation resistance in Fig. 17 and Fig. 18 showing that the reduction in grain size is reducing the propagation rates of microcracks in Al-alloys [8] as well as in a + b Ti-alloys. The pancake-shaped grain structure in Al-alloys results in
anisotropic mechanical properties depending on the testing direction. The evaluation in Table 2 is valid for the ‘good’ testing direction, that means the stress axis is parallel to the longitudinal direction of the pancakes (L-direction). Soft zones in the material (GB a-layers in a+b Ti-alloys and PFZ in Al-alloys) have a deleterious effect on some of the important mechanical properties (Table 2). The negative effect is increasing with the strength difference between the hard matrix and the soft zones. This is illustrated in Fig. 19 in which the tensile elongation is plotted as a function of cooling rate from the b phase field (fully lamellar structures) or from the bi-modal recrystallization temperature. With increasing cooling rate the a-lamellae size is decreased and the yield stress is increased especially in the faster cooling regime above 100°C min − 1 in which an increasing amount of martensitic plates are generated. The slip length is therefore decreased with increasing cooling rate and this should increase the tensile ductility. But, as can be seen from Fig. 19 for the fully lamellar condition with a large b grain size of 600 mm, the tensile ductility is passing through a maximum with increasing cooling rate and beyond the maximum the fracture mechanism is shifting to a ductile fracture along the soft GB a-layers. With decreasing b grain size (Fig. 19, 100 mm values) the effective length of the GB a-layers is decreased resulting in an improvement in ductility and less ‘grain hondary’ failure [9]. The bi-modal structure with a small b grain size of 25 mm shows the predicted steady increase in ductility (reduction in area values), the observed slight decrease in tensile elongation at fast cooling rates (Fig. 19) is due to a decrease in strain hardening exponent and earlier on-set of necking.
Table 5 Influence of grain size and shape on tensile ductility and fracture toughness for b Ti-alloys and overaged Al-alloys b Ti-alloys sy =1200 (MPa)
RA (%)
KIC (MPa m−1/2)
RA (%)
KIC (MPa m−1/2)
Overaged Al-alloys sy =480 (MPa)
Large GS Small GS (bi-modal) Necklace (L) Necklace (ST)
10 35 16 10
52 37 68 37
10 35 35 16
40 33 47 31
Large GS Small GS Pancake (L) Pancake (ST)
Fig. 24. b CEZ, crack front profiles of fracture toughness sepcimens, LM. (a) Large grain size (lamellar). (b) Small grain size (bi-modal). (c) Necklace (L). (d) Necklace (ST).
G. Lu¨tjering / Materials Science and Engineering A263 (1999) 117–126
125
Fig. 28. Al-Alloy X-7075, overaged, crack nucleation, LM. (a) Large grain size 220 mm. (b) Small grain size 30 mm. Fig. 25. b-CEZ, influence of grain size on HCF (R = −1).
the HCF strength (Fig. 21). The two S–N curves in Fig. 21 demonstrate the beneficial effect of reducing the Fe and Si content from the older 7075 alloy to the newer 7475 alloy.
5.2. Comparison of b Ti-alloys and o6eraged Al-alloys
Fig. 26. Al-Alloy X-7075, overaged, influence of grain size on HCF (R = −1).
The deliberately added dispersoids have, besides the indirect influence on properties through the necklace grain shape, also a direct influence on properties of underaged Al-alloys because they reduce the formation of intense slip bands improving ductility and HCF strength. Unfortunately, the crack propagation properties are negatively influenced by the dispersoids because they nucleate voids in front of the crack tip. The large Fe and Si containing inclusions serve as crack nucleation sites in commercial Al-alloys (Fig. 20) lowering
As can be seen from Table 3, the agreement in microstructural parameters (except for the dispersoids and inclusions) and in the structure/property correlations is even better for the comparison of b Ti-alloys and overaged Al-alloys as it was for a+ b Ti-alloys and underaged Al-alloys. With increasing overaging (reduction in yield stress), all crack propagation properties are improved because the ductility is increased and large non-shearable plates hinder crack propagation. This is illustrated in Fig. 22 and Fig. 23 showing da/dN DK curves of macrocracks. In the case of b Ti-alloys the a plate size in b-CEZ was varied drastically from about 0.3 mm (fine lamellar) to about 50 mm (coarse lamellar) and the resulting effect on threshold was 2 MPa m1/2 (Fig. 22) [10]. As mentioned already earlier, the effect of PFZ on properties is more pronounced for overaged Al-alloys, but the effect is qualitatively the same as for underaged Al-alloys (Table 2). Table 5 demonstrates the influence of grain size and grain shape on tensile ductility and fracture toughness for b Ti-alloys and
Fig. 27. b-CEZ, crack nucleation, LM. (a) Large grain size (lamellar). (b) Small grain size (bi-modal).
126
G. Lu¨tjering / Materials Science and Engineering A263 (1999) 117–126
overaged Al-alloys containing GB a-layers and PFZ, respectively. By reducing the grain size the ductility is increased as discussed already on the example of GB a-layers in a +b Ti-alloys. The same is true for changing the grain shape to necklace or pancake and testing in L-direction, but for these anisotropic grain shapes the ductility is reduced drastically for tests in the ST-direction (Table 5). It should be mentioned, that aligning the soft zones under 45° to the stress axis results in even lower ductilities. The fracture toughness shows a quite different dependence (Table 5) because the crack front profile has to be taken into account for the macrocracks (Fig. 24). The conditions with a small grain size and with a necklace/pancake structure tested in ST-direction exhibit a very flat crack front geometry and low fracture toughness values (Table 5) whereas the other two conditions with a rough crack front profile exhibit higher fracture toughness values (Table 5). The last example (Figs. 25 – 28) demonstrates the effect of grain size on HCF (crack nucleation resistance) for b Ti-alloys and overaged Al-alloys containing GB a-layers and PFZ, respectively. In both cases the crack nucleation resistance is increased with decreasing grain size (Fig. 25 and Fig. 26) because the effective length of the soft zones is decreased. For the conditions with large grain size the cracks are nucleated at GB a-layers (Fig. 27a) and at PFZ (Fig. 28a). For the conditions with small grain size the overaged Al-alloy shows still crack nucleation at the PFZ (Fig. 28b), whereas for the bi-modal structure of the b Ti-alloy cracks are nucleated at large a plates having the same length as the b grain size (Fig. 27b), representing therefore easier crack nucleation sites than the GB a-layers. The direct influence of the dispersoids on properties which was present for underaged Al-alloys (Table 2) disappears for overaged Al-alloys for those mechanical properties which are solely determined by the PFZ, namely oF, HCF, da/dN of microcracks, and KIC (Table 3). Naturally, the positive indirect influence of the dispersoids on those properties through the pancake grain structure is still present.
6. Conclusions Similarities in the microstructure/property relation-
.
ships exist between a+ b Ti-alloys and underaged Alalloys and between b Ti-alloys and overaged Al-alloys. Differences are caused by the presence of dispersoids and inclusions in Al-alloys and by the occurrence of the alloying element partitioning effect in bi-modal structures of a+ b Ti-alloys. The latter effect can be reduced by an intermediate heat treatment at 800–850°C. From the two possibilities to reduce the negative effects of soft zones along grain boundaries (GB a-layers in Ti-alloys and PFZ in Al-alloys), namely small equiaxed grain sizes or elongated grain structures (pancakes), the processing route for achieving small grain sizes is much easier in Ti-alloys (bi-modal microstructure), whereas in Al-alloys the generation of a pancake grain structure by processing is very easy due to the presence of the dispersoids inhibiting recrystallization.
Acknowledgements The author would like to express his appreciation to Professor J. Albrecht, Professor A. Gysler, Dr J.-O. Peters, and Professor L. Wagner (now at TU Cottbus) for contributing with their research activities at the TU Hamburg-Harburg significantly to the content of this paper.
References [1] M. Peters, G. Lu¨tjering, G. Ziegler, Z. Metallkde 74 (1983) 274. [2] J.O. Peters, G. Lu¨tjering, M. Koren, H. Puschnik, R.R. Boyer, Mat. Sci. Eng., A 213 (1996) 71. [3] J. Waldman, H. Sulinski, H. Markus, Met. Trans. 5 (1974) 573. [4] G. Lu¨tjering, A. Gysler, 5th Intern. Conf. on Titanium, DGM (1985) 2065. [5] G. Lu¨tjering, A. Gysler, L. Wagner, 6th World Conf. on Titanium, les e´ditions de physique (1988) 71. [6] G. Lu¨tjering, A. Gysler, Aluminum Alloys, EMAS (1986) 1547. [7] G. Lu¨tjering, A. Gysler, Proc. 8th Intern. Light Metals Congress, Leoben, Austria (1987) 649. [8] L. Wagner, G. Lu¨tjering, Fatigue ‘87, EMAS (1987) 1819. [9] G. Lu¨tjering, J. Albrecht, O.M. Invasishin, 8th World Conf. on Titanium, The University Press, Cambridge, UK, 1996, p. 1163. [10] J.O. Peters, C. Sauer, G. Lu¨tjering, Fatigue Behavior of Titanium Alloys, TMS Fall Meeting ‘98, TMS (to be published).
Property optimization through microstructural control in titanium and aluminum alloys G. Lu¨tjering * Technical Uni6ersity Hamburg-Harburg, 21071 Hamburg, Germany
Abstract In this paper the microstructure/property relationships of Ti-alloys (a + b and b alloys) and high strength, age-hardened Al-alloys are discussed. Common features in these microstructure/property relationships can be found for a+ b Ti-alloys and Al-alloys in the underaged condition as well as for b Ti-alloys and Al-alloys in the overaged condition. Basic differences are the presence of dispersoids and large Fe and Si containing inclusions in Al-alloys and the alloying element partitioning effect in bi-modal structures of a +b Ti-alloys. © 1999 Elsevier Science S.A. All rights reserved. Keywords: Al-alloys; Fe; Si; Ti-alloys; Microstructure; Mechanical properties
1. Introduction
2. Microstructure of a+ b Ti-alloys
Titanium alloys (a + b and b alloys) and aluminum alloys are widely used in aerospace applications because of their low density and high strength values. Typical values for some important mechanical properties are shown in Table 1 for Ti – 6Al – 4V (a + b Ti-alloy), b-CEZ (b Ti-alloy) and the high strength, age-hardened Al-alloy 7475. It can be seen that the a + b Ti-alloy Ti –6Al–4V shows the best resistance to crack propagation whereas the b Ti-alloy b-CEZ exhibits the highest tensile and fatigue strength. Comparing these values to the Al-alloy 7475, the difference in density (4.5 g cm − 3 for Ti-alloys and 2.7 g cm − 3 for Al-alloys, a factor of about 1.7) should be taken into account. For all three types of alloys mechanical properties are optimized to a large extent through microstructural control by appropriate processing and heat treatment. The present paper tries to outline common features in such microstructural control and in the microstructure/property relationships as well as basic differences. Common guide-lines for these microstructure/property relationships might be useful for further improvements in mechanical properties for both Ti-alloys as well as for Al-alloys.
The main microstructural features of fully lamellar structures in a+ b Ti-alloys (Fig. 1) are the relatively large size of the equiaxed b grains (about 600 mm), the presence of continuous a-layers at b grain boundaries, the a-colony size and the size of the individual a-lamellae. Since the recrystallization kinetics in the b phase field are so fast, only the cooling rate from the b phase field has a major influence on microstructure in the whole processing route. With increasing cooling rate the size of the a-colonies as well as the size of the individual a-lamellae are reduced, at extreme fast cooling rates down to the dimensions of a single martensitic plate. For nearly all a+ b Ti-alloys the a phase is age-hardened by the precipitation of very small (nanometer sized), coherent Ti3Al particles (a2) during the final low temperature heat treatment, such as 24 h 500°C (Ti– 6Al–4V), 8 h 590°C (Ti-6242), 2 h 700°C (IMI 834). The volume fraction of these a2 particles depends also on the oxygen content of the alloy, because oxygen is known to stabilize the a2 phase. In contrast to the fully lamellar structures (Fig. 1), a so-called bi-modal (or duplex) type of microstructure (Fig. 2) can be obtained in a+ b Ti-alloys by thermomechanical processing (Fig. 3). In this case the lamellar microstructure is deformed in the a+ b phase field and then recrystallized in the a+b phase field leading to a mixture of equiaxed a and
* Tel.: +49-40-7718-3036; fax: +49-40-7718-2325. E-mail address: [email protected] (G. Lu¨tjering)
0921-5093/99/$ - see front matter © 1999 Elsevier Science S.A. All rights reserved. PII: S 0 9 2 1 - 5 0 9 3 ( 9 8 ) 0 1 1 6 9 - 1
G. Lu¨tjering / Materials Science and Engineering A263 (1999) 117–126
118
Table 1 Typical values for mechanical properties at room temperature Alloy
Ti–6Al–4V (bi-modal) b-CEZ (bi-modal) Al-7475 (pancakes)
E (GPa)
sy (MPa)
T.E. (%)
HCF R =−1 (MPa)
DK at 10-9 (m cy.−1)
KIC (MPa m−1/2)
Microcracks R=−1 (MPa m−1/2)
Macrocracks R =0.1 (MPa m−1/2)
115
950
22
500
3
8
60
120
1200
13
625
2
4.5
37
70
480
14
200
1
3.5
45
b grains [1]. Upon cooling from the bi-modal recrystallization temperature the b grains are transformed again to a lamellar structure. The main microstructural difference to fully lamellar structures is the small b grain size (about 20–50 mm) determined by the distance of the primary a grains (ap). This small b grain size of bimodal structures limits the maximum a-colony size and the maximum length of a-lamellae as well as the effective length of the GB a-layers. The most important parameter of the processing route is the correct choice of the recrystallization or bi-modal annealing temperature because it determines the volume fraction of ap. The aging response of bi-modal microstructures is somewhat different from fully lamellar microstructures because of the alloying element partitioning effect taking place during the bi-modal recrystallization treatment and leading to an enrichment in aluminum and oxygen in the ap phase. As a consequence, the lamellar grains in the bi-modal structure are softer (less a2 particles) as compared to the fully lamellar structure.
temperature. Instead, the a phase is precipitated in form of incoherent a-plates within the b grains upon annealing or aging in the a+ b phase field. The main difference between the various microstructures in b Ti-alloys is the size and shape of the b grains (Figs. 4–6). The size of the equiaxed b grains, resulting from recrystallization in the b phase field, is somewhat smaller (about 400 mm) than in a+ b Ti-alloys because of the 150–200°C lower b transus. Upon cooling from the b phase field, extensive GB a-layers are formed (Fig. 4). The heat treatment in the a+ b phase field consists usually of two steps. First, the annealing is performed at relatively high tremperatures leading to
3. Microstructure of b Ti-alloys The characteristic feature of b Ti-alloys is that the content of b stabilizing elements is high enough to suppress martensitic transformation as well as the formation of a a-colony structure upon cooling to room
Fig. 2. a+ b Ti-Alloys, bi-modal structure.
Fig. 1. a+ b Ti-Alloys, lamellar structure.
Fig. 3. a +b Ti-Alloys, processing route for bi-modal structures (schematically).
G. Lu¨tjering / Materials Science and Engineering A263 (1999) 117–126
Fig. 4. b Ti-Alloys, lamellar structure.
119
Fig. 8. b Ti-Alloys, through-transus processing route for necklace structures (schematically).
Fig. 5. b-Ti-Alloys, bi-modal structure.
the formation of large a-plates visible in the LM-picture of Fig. 4. The second step is an aging treatment at lower temperatures (around 600°C) precipitating fine aplatelets shown in the TEM-picture of Fig. 7. The
Fig. 6. b Ti-Alloys, necklace structure.
Fig. 7. b Ti-Alloys, secondary a-platelets.
Fig. 9. Al-Alloys, pancake grain structure.
resulting microstructure of this coarse-grained b recrystallized condition (Fig. 4) is commonly called ‘lamellar’. To avoid the negative influence of the extensive GB a-layers on mechanical properties, two different approaches have been taken for b Ti-alloys [2]. One approach is to obtain a bi-modal type microstructure (Fig. 5) similar to the bi-modal microstructure in a+ b Ti-alloys. The other approach is to create a so-called necklace type microstructure (Fig. 6). The thermomechanical processing route to obtain bi-modal microstructures is similar to the route for a+ b Ti-alloys shown in Fig. 3, except that after the bi-modal recrystallization treatment the above described two-stage heat treatment is used for b Ti-alloys to form large a-plates and small a-platelets. To create a necklace type microstructure (Fig. 6) a so-called ‘through-transus’ deformation process has to be applied (Fig. 8). After homogenization the material is not cooled to room temperature but directly deformed starting in the b phase field and continued into the a+b phase field until a forms at the deformed b grain boundaries as round particles stabilizing the elongated or pancake shaped b grain structure (Fig. 6). It is important that the deformation process is finished before a starts to precipitate also within the interior of the deformed b grains as round particles. Upon cooling, the b grains recover and the final two-step heat treatment forms the desired mixture of large a-plates and small a-platelets (Fig. 7) also in this necklace type of microstructure.
120
G. Lu¨tjering / Materials Science and Engineering A263 (1999) 117–126
4. Microstructure of high strength Al-alloys High strength Al-alloys are age-hardened alloys based on the Al– Cu system (2XXX series) or on the Al–Zn system (7XXX series). Particles of the equilibrium phase, for example Al2Cu (u), are incoherent with the fcc a phase, but upon aging at temperatures below about 200°C small metastable precipitates (size 2 –20 nm) are formed (first coherent u¦ and then semicoherent u%) leading to the characteristic age-hardening curve in which the yield stress is passing through a maximum with aging time. In the underaged condition (before the maximum), coherent u¦ particles are present which can be sheared by the moving dislocations leading to the formation of intense slip bands. In the overaged condition (after the maximum), semi-coherent u% particles are present which cannot be sheared that easily and are mostly by-passed by the moving dislocations (Orowan–mechanism) creating dislocation loops around the particles and a more homogeneous slip distribution. In addition, two other kinds of particles are present in commercial Al-alloys: small incoherent dispersoids (size about 0.2 mm) containing alloying elements such as Mn, Cr, Ti forming intermetallic compounds (Al6Mn, Al3Cr, Al3Ti) and large inclusions (size about 5–10 mm) containing the impurity elements Fe and Si. Both, the dispersoids and the large inclusions are stable up to the melting point. The resulting
Fig. 12. Al-Alloys, dispersoids and subgrains.
Fig. 13. Al-Alloys, processing route for pancake structure (schematically).
microstructural features of high strength Al-alloys are illustrated in Figs. 9–12. As a consequence of the deliberately added, homogeneously distributed dispersoids (Fig. 12) inhibiting recrystallization, high strength Al-alloys exhibit usually an unrecrystallized pancakeshaped grain structure (Fig. 9). The processing route (Fig. 13) consists of homogenization in the single a phase field (at around 500°C) followed by deformation in the temperature region of about 300–400°C (two phase region). During the subsequent solution treatment in the a phase field all precipitated equilibrium phase is re-solutioned but the material is not recrystallizing, only recovery into small subgrains takes place (Fig. 12). The last step in the processing route is the aging treatment to precipitate the coherent or semi-co-
Fig. 10. Al-Alloys, recrystallized grain structure.
Fig. 11. Al-Alloys, PFZ along grain boundaries.
Fig. 14. Al-Alloys, processing route for recrystallized structure (schematically).
G. Lu¨tjering / Materials Science and Engineering A263 (1999) 117–126
121
Table 2 Influence of microstructural parameters on mechanical properties of a+b Ti-alloys and underaged Al-alloys a+b Titanium alloys
sy
Aging (a2), oxygen Bi-modal structure GB a-layers Small a-colonies, alamellae
oF
HCF
Micro-cracks da/dn
Macrocracks
DKth R = 0.7
KIC
DKth R = 0.1
Creep strength 0.2%
Underaged aluminum alloys
+ + O +
− + − +
+ − + − +
− + − +
− − O −
− − − −
+ − O −
+ − O O −
Aging (coh. prec.) Pancakes, small GS PFZ
O
+
+
−
−
−
−
+
O
−
−
O
O
−
O
O
Dispersoids (Mn, Cr, Ti) Inclusions (Fe,Si)
Table 3 Influence of microstructural parameters on mechanical properties of b Ti-alloys and overaged Al-alloys b Titanium alloys
Overaging (asec) GB a-layers Bi-modal structure Necklace structure
sy
oF
HCF
Micro-cracks da/dn
Macrocracks
Creep strength 0.2% Overaged aluminum alloys
DKth R = 0.7
KIC
DKth R = 0.1
− O O
+ − +
− − +
+ − +
+ O −
+ − −
+ O −
− O O
Overaging (semi-coh. prec.) PFZ Small GS
O
+
+
+
−
+
−
O
Pancakes
O O
O −
O −
O O
− O
O −
− O
+ O
Dispersoids (Mn, Cr, Ti) Inclusions (Fe,Si)
herent particles. Unfortunately, particles of the incoherent equilibrium phase, for example Al2Cu (u), are formed during the aging treatment at the grain boundaries creating soft, precipitate-free-zones (PFZ) along the grain boundaries (Fig. 11). The formation of the PFZ is starting already in the underaged stage before the maximum in yield stress is reached but the negative effects of the PFZ are much more pronounced for the overaged condition. In order to obtain recrystallized microstructures with small, nearly equiaxed a grains (Fig. 10), a fairly complicated processing route has to be applied (Fig. 14) [3]. After homogenization the material is annealed in the two phase region to precipitate the equilibrium phase as very coarse particles. The subsequent deformation process is performed at relatively low temperatures followed by recrystallization in the two phase region at temperatures close to the phase boundary. Afterwards a short solution treatment is performed to re-solution the equiblibrium phase followed by the final aging treatment. It should be pointed out that this complicated processing route is not very practical, another way is to reduce the content of the dispersoid-forming elements and to deform the material in the normal processing route (Fig. 13) at as
low as possible temperatures. The resulting grain structure is not as uniform as shown in Fig. 10, the grain size is in most cases fairly inhomogeneous due to the reduced volume fraction of dispersoids.
5. Correlation between microstructure and mechanical properties In this correlation between microstructure and mechanical properties the a+ b Ti-alloys can be compared to Al-alloys in the underaged condition (Table 2), because both materials contain coherent, shearable precipitates, whereas the b Ti-alloys should be compared to Al-alloys in the overaged condition (Table 3), because both materials contain non-shearable precipitates (incoherent a-platelets in case of b Ti-alloys and semicoherent precipitates in case of overaged Al-alloys). In Table 2 and Table 3 the relevant microstructural parameters are listed together with important mechanical properties and the experimentally observed influence is qualitatively indicated by + (positive), − (negative), or O (no influence). Within the scope of this paper, it is not intended to discuss all the individual relationships
122
G. Lu¨tjering / Materials Science and Engineering A263 (1999) 117–126
Table 4 Influence of aging on fracture toughness Ti-alloy Ti–8.6Al
Fracture toughness (MPa m−1/2)
10 h 500°C 260 h 500°C
75 70
[4 –7] in detail. In the following only a rough overview will be given with some examples.
5.1. Comparison of a + b Ti-alloys and underaged Al-alloys Increasing the aging degree in underaged Al-alloys and in a +b Ti-alloys as well as the oxygen content in the latter case results in the same qualitative influence on all mechanical properties listed in Table 2. Increasing the degree of age-hardening by coherent, shearable
Al-alloy X-7075 55 40
24 h 100°C 20 h 160°C
particles will increase the intensity of slip band formation reducing the ductility. The negative influence on fatigue crack nucleation is usually overcompensated by the increase in yield stress. The fatigue crack propagation of small, naturally initiated surface cracks (microcracks) which propagate in this stage in single slip bands (flat crack front) is increased with the increasing intensity of the slip bands. Large macrocracks (CT type of
Fig. 17. Ti-Alloy IMI 834, microcrack propagation (R= − 1). Fig. 15. Ti-Alloy IMI 834, HCF at RT, R = −1, fully lamellar and bi-modal microstructures.
Fig. 16. Al-Alloy X-7075, underaged, HCF at RT, R = −1, influence of grain size.
Fig. 18. Al-Alloy X-7075, underaged, microcrack propagation (R = −1).
G. Lu¨tjering / Materials Science and Engineering A263 (1999) 117–126
Fig. 19. Ti – 6Al – 4V, effect of cooling rate on tensile elongation.
123
erties as changing from fully lamellar structures to bi-modal structures for a+ b Ti-alloys, except for two properties, HCF and creep strength, (Table 2). In these two cases, Table 2 shows in the left column the dependence for bi-modal structures of a+ b Ti-alloys and in the right column the dependence for Al-alloys. This different behavior is demonstrated for the HCF strength by comparing Fig. 15 and Fig. 16 and can be attributed to the alloying element partitioning effect in bi-modal structures of a+ b Ti-alloys leading to crack nucleation in the lamellar regions of the bi-modal structure at lower stress amplitudes than in fully lamellar structures. This alloy element partitioning effect in bimodal structures of a+ b Ti-alloys can also explain the difference between Al-alloys and a+ b Ti-alloys in case of creep strength. It should be pointed out that the negative influence of the alloy element partitioning effect in a+b Ti-alloys can be eliminated by introduc-
Fig. 20. Al-7475, fatigue crack nucleation at large inclusions.
specimens), in the absence of crack closure (high Rvalue, KIC), exhibit lower DKth and KIC values with increasing aging because the crack front profile is dominated by the grain size and is not significantly changed by the intensity of slip bands. The decrease in fracture toughness with increasing degree of age-hardening is illustrated in Table 4 for Ti- and Al-alloys. The fatigue crack propagation threshold value of macrocracks at low R-values is usually increased with increasing aging (Table 2), because the crack closure contribution, increasing with increasing shear displacement at the crack tip (mode II), usually overcompensates the easier fatigue crack propagation with increasing aging (see microcracks). The creep resistance listed in Table 2 (opl. + 0.2%, primary creep) is valid for the dislocationcreep regime (up to 200°C for Al-alloys and 600°C for Ti-alloys). This creep resistance is increasing with agehardening degree (increasing yield stress). Reducing the a grain size or changing to pancake-shaped grains for Al-alloys show the same tendency on mechanical prop-
Fig. 21. Al-Alloys, influence of Fe and Si content on HCF (R = −1).
Fig. 22. b-CEZ, effect of size of a-plates on da/dN– DK curves.
G. Lu¨tjering / Materials Science and Engineering A263 (1999) 117–126
124
Fig. 23. Al-Alloy X-7075, effect of overaging on da/dN– DK curves.
ing in the processing route (Fig. 3) an intermediate heat treatment in the temperature region of 800 – 850°C between the bi-modal recrystallization treatment and the final aging treatment. The alloy element partitioning effect does not play a significant role for all other mechanical properties listed in Table 2. This is illustrated by the example of microcrack propagation resistance in Fig. 17 and Fig. 18 showing that the reduction in grain size is reducing the propagation rates of microcracks in Al-alloys [8] as well as in a + b Ti-alloys. The pancake-shaped grain structure in Al-alloys results in
anisotropic mechanical properties depending on the testing direction. The evaluation in Table 2 is valid for the ‘good’ testing direction, that means the stress axis is parallel to the longitudinal direction of the pancakes (L-direction). Soft zones in the material (GB a-layers in a+b Ti-alloys and PFZ in Al-alloys) have a deleterious effect on some of the important mechanical properties (Table 2). The negative effect is increasing with the strength difference between the hard matrix and the soft zones. This is illustrated in Fig. 19 in which the tensile elongation is plotted as a function of cooling rate from the b phase field (fully lamellar structures) or from the bi-modal recrystallization temperature. With increasing cooling rate the a-lamellae size is decreased and the yield stress is increased especially in the faster cooling regime above 100°C min − 1 in which an increasing amount of martensitic plates are generated. The slip length is therefore decreased with increasing cooling rate and this should increase the tensile ductility. But, as can be seen from Fig. 19 for the fully lamellar condition with a large b grain size of 600 mm, the tensile ductility is passing through a maximum with increasing cooling rate and beyond the maximum the fracture mechanism is shifting to a ductile fracture along the soft GB a-layers. With decreasing b grain size (Fig. 19, 100 mm values) the effective length of the GB a-layers is decreased resulting in an improvement in ductility and less ‘grain hondary’ failure [9]. The bi-modal structure with a small b grain size of 25 mm shows the predicted steady increase in ductility (reduction in area values), the observed slight decrease in tensile elongation at fast cooling rates (Fig. 19) is due to a decrease in strain hardening exponent and earlier on-set of necking.
Table 5 Influence of grain size and shape on tensile ductility and fracture toughness for b Ti-alloys and overaged Al-alloys b Ti-alloys sy =1200 (MPa)
RA (%)
KIC (MPa m−1/2)
RA (%)
KIC (MPa m−1/2)
Overaged Al-alloys sy =480 (MPa)
Large GS Small GS (bi-modal) Necklace (L) Necklace (ST)
10 35 16 10
52 37 68 37
10 35 35 16
40 33 47 31
Large GS Small GS Pancake (L) Pancake (ST)
Fig. 24. b CEZ, crack front profiles of fracture toughness sepcimens, LM. (a) Large grain size (lamellar). (b) Small grain size (bi-modal). (c) Necklace (L). (d) Necklace (ST).
G. Lu¨tjering / Materials Science and Engineering A263 (1999) 117–126
125
Fig. 28. Al-Alloy X-7075, overaged, crack nucleation, LM. (a) Large grain size 220 mm. (b) Small grain size 30 mm. Fig. 25. b-CEZ, influence of grain size on HCF (R = −1).
the HCF strength (Fig. 21). The two S–N curves in Fig. 21 demonstrate the beneficial effect of reducing the Fe and Si content from the older 7075 alloy to the newer 7475 alloy.
5.2. Comparison of b Ti-alloys and o6eraged Al-alloys
Fig. 26. Al-Alloy X-7075, overaged, influence of grain size on HCF (R = −1).
The deliberately added dispersoids have, besides the indirect influence on properties through the necklace grain shape, also a direct influence on properties of underaged Al-alloys because they reduce the formation of intense slip bands improving ductility and HCF strength. Unfortunately, the crack propagation properties are negatively influenced by the dispersoids because they nucleate voids in front of the crack tip. The large Fe and Si containing inclusions serve as crack nucleation sites in commercial Al-alloys (Fig. 20) lowering
As can be seen from Table 3, the agreement in microstructural parameters (except for the dispersoids and inclusions) and in the structure/property correlations is even better for the comparison of b Ti-alloys and overaged Al-alloys as it was for a+ b Ti-alloys and underaged Al-alloys. With increasing overaging (reduction in yield stress), all crack propagation properties are improved because the ductility is increased and large non-shearable plates hinder crack propagation. This is illustrated in Fig. 22 and Fig. 23 showing da/dN DK curves of macrocracks. In the case of b Ti-alloys the a plate size in b-CEZ was varied drastically from about 0.3 mm (fine lamellar) to about 50 mm (coarse lamellar) and the resulting effect on threshold was 2 MPa m1/2 (Fig. 22) [10]. As mentioned already earlier, the effect of PFZ on properties is more pronounced for overaged Al-alloys, but the effect is qualitatively the same as for underaged Al-alloys (Table 2). Table 5 demonstrates the influence of grain size and grain shape on tensile ductility and fracture toughness for b Ti-alloys and
Fig. 27. b-CEZ, crack nucleation, LM. (a) Large grain size (lamellar). (b) Small grain size (bi-modal).
126
G. Lu¨tjering / Materials Science and Engineering A263 (1999) 117–126
overaged Al-alloys containing GB a-layers and PFZ, respectively. By reducing the grain size the ductility is increased as discussed already on the example of GB a-layers in a +b Ti-alloys. The same is true for changing the grain shape to necklace or pancake and testing in L-direction, but for these anisotropic grain shapes the ductility is reduced drastically for tests in the ST-direction (Table 5). It should be mentioned, that aligning the soft zones under 45° to the stress axis results in even lower ductilities. The fracture toughness shows a quite different dependence (Table 5) because the crack front profile has to be taken into account for the macrocracks (Fig. 24). The conditions with a small grain size and with a necklace/pancake structure tested in ST-direction exhibit a very flat crack front geometry and low fracture toughness values (Table 5) whereas the other two conditions with a rough crack front profile exhibit higher fracture toughness values (Table 5). The last example (Figs. 25 – 28) demonstrates the effect of grain size on HCF (crack nucleation resistance) for b Ti-alloys and overaged Al-alloys containing GB a-layers and PFZ, respectively. In both cases the crack nucleation resistance is increased with decreasing grain size (Fig. 25 and Fig. 26) because the effective length of the soft zones is decreased. For the conditions with large grain size the cracks are nucleated at GB a-layers (Fig. 27a) and at PFZ (Fig. 28a). For the conditions with small grain size the overaged Al-alloy shows still crack nucleation at the PFZ (Fig. 28b), whereas for the bi-modal structure of the b Ti-alloy cracks are nucleated at large a plates having the same length as the b grain size (Fig. 27b), representing therefore easier crack nucleation sites than the GB a-layers. The direct influence of the dispersoids on properties which was present for underaged Al-alloys (Table 2) disappears for overaged Al-alloys for those mechanical properties which are solely determined by the PFZ, namely oF, HCF, da/dN of microcracks, and KIC (Table 3). Naturally, the positive indirect influence of the dispersoids on those properties through the pancake grain structure is still present.
6. Conclusions Similarities in the microstructure/property relation-
.
ships exist between a+ b Ti-alloys and underaged Alalloys and between b Ti-alloys and overaged Al-alloys. Differences are caused by the presence of dispersoids and inclusions in Al-alloys and by the occurrence of the alloying element partitioning effect in bi-modal structures of a+ b Ti-alloys. The latter effect can be reduced by an intermediate heat treatment at 800–850°C. From the two possibilities to reduce the negative effects of soft zones along grain boundaries (GB a-layers in Ti-alloys and PFZ in Al-alloys), namely small equiaxed grain sizes or elongated grain structures (pancakes), the processing route for achieving small grain sizes is much easier in Ti-alloys (bi-modal microstructure), whereas in Al-alloys the generation of a pancake grain structure by processing is very easy due to the presence of the dispersoids inhibiting recrystallization.
Acknowledgements The author would like to express his appreciation to Professor J. Albrecht, Professor A. Gysler, Dr J.-O. Peters, and Professor L. Wagner (now at TU Cottbus) for contributing with their research activities at the TU Hamburg-Harburg significantly to the content of this paper.
References [1] M. Peters, G. Lu¨tjering, G. Ziegler, Z. Metallkde 74 (1983) 274. [2] J.O. Peters, G. Lu¨tjering, M. Koren, H. Puschnik, R.R. Boyer, Mat. Sci. Eng., A 213 (1996) 71. [3] J. Waldman, H. Sulinski, H. Markus, Met. Trans. 5 (1974) 573. [4] G. Lu¨tjering, A. Gysler, 5th Intern. Conf. on Titanium, DGM (1985) 2065. [5] G. Lu¨tjering, A. Gysler, L. Wagner, 6th World Conf. on Titanium, les e´ditions de physique (1988) 71. [6] G. Lu¨tjering, A. Gysler, Aluminum Alloys, EMAS (1986) 1547. [7] G. Lu¨tjering, A. Gysler, Proc. 8th Intern. Light Metals Congress, Leoben, Austria (1987) 649. [8] L. Wagner, G. Lu¨tjering, Fatigue ‘87, EMAS (1987) 1819. [9] G. Lu¨tjering, J. Albrecht, O.M. Invasishin, 8th World Conf. on Titanium, The University Press, Cambridge, UK, 1996, p. 1163. [10] J.O. Peters, C. Sauer, G. Lu¨tjering, Fatigue Behavior of Titanium Alloys, TMS Fall Meeting ‘98, TMS (to be published).