Mechanisms influencing the cryogenic fracture-toughness behavior of aluminum-lithium alloys

$Page 1: Mechanisms influencing the cryogenic fracture-toughness behavior of aluminum-lithium alloys$
Acta metall, mater. Vol. 38, No. 11, pp. 2309-2326, 1990 0956-7151/90 $3.00 + 0.00 Printed in Great Britain. All rights reserved Copyright © 1990 Pergamon Press plc

MECHANISMS INFLUENCING THE CRYOGENIC FRACTURE-TOUGHNESS BEHAVIOR OF

ALUMINUM-LITHIUM ALLOYS

K. T. VENKATESWARA RAO and R. O. RITCHIE Center for Advanced Materials, Lawrence Berkeley Laboratory and Department of Materials Science and

Mineral Engineering University of California, Berkeley, CA 94720, U.S.A.

(Received 23 December 1989)

Abstract--Cryogenic strength-toughness relationships for advanced aluminum-lithium alloys 2090, 8090, 8091 and 2091 are examined as a function of microstructure, plate orientation and wrought-product form (plate vs sheet), with specific emphasis on the underlying micro-mechanisms associated with crack advance. It is found that, with decrease in temperature from 298 K to 77 and 4 K, strength, tensile elongation and strain-hardening exponent are increased for all alloy chemistries, microstructures and product forms; however, the longitudinal (L-T, T-L) fracture toughness may increase or decrease depending upon the prevailing microscopic mechanism (microvoid coalescence vs transgranular shear) and macroscopic mode (plane strain vs plane stress) of fracture. In general, alloy microstructures that exhibit changes in either the fracture mechanism or mode at low temperatures show a decrease in L-T toughness. Conversely, when the fracture mechanism is unchanged between ambient and 4 K, observed variations in toughness with temperature are a strong function of the degree of local stress-triaxiality that develops at the crack tip. In very thin sheets, where the fracture mode remains one of plane stress ("slant" fracture), the elevation in toughness at low temperatures is associated with the concurrent increase in tensile strength and ductility; conversely, in thick plate, the increased occurrence of through-thickness delaminations (due to the weak short-transverse properties) at low temperatures locally promotes plane-stress conditions, thereby enhancing toughness by relaxing triaxial constraint. In sheets of intermediate thickness, however, the absence of such through-thickness delaminations permits the expected transition from plane-stress to plane-strain conditions, with the result that the toughness now decreases with reduction in temperature.

R r s u m r - ~ n examine les relations entre rrsistance mrcanique h basse temprrature et tenacit6 pour des alliages avancrs aluminium-lithium 2090, 8090, 8091 et 2091 en fonction de la microstructure, de l'orientation de la plaque et de la forme du produit forg6 (plaque ou trle), avec un accent particulier sur les micromrcanismes sous-jacents associrs ~ la propagation des fissures. On trouve que, lorsque la temperature drcroit de 298 ~ 77 et/ t 4 K, la r~sistance mrcanique, l'61ongation en traction et l'exposant d'rcrouissage augrnentent quelles que soient la composition chimique de l'alliage, sa microstructure et sa forme; cependant, la tenacit6 en fracture longitudinale (L-T, T-L) peut augmenter ou diminuer suivant le mrcanisme de rupture dominant; microscopique (coalescence de microcavitrs en fonction du cisaillement transgranulaire) ou macroscopique (drformation plane en fonction de la contrainte plane). En grnrral, les microstructures d'alliage qui montrent des variations soit du mrcanisme, soit du mode de rupture ~i basses temprratures, rrvrlent une drcroissance de la trnacit6 L-T. Inversement, lorsque le mrcanisme de rupture est inchang6 entre la temprrature ambiante et 4K, les variations observres de la tenacit6 en fonction de la temprrature drpendent fortement du degr6 de triaxialit6 de la contrainte locale qui se drveloppe l'extrrmit6 de la fissure. Dans les films trrs minces, oti le mode de rupture reste du type contrainte plane (rupture "inclinre"), rrlrvation de la tenacit6 aux basses temprratures est associre fi l'augmentation simultanre de la rrsistance mrcanique en traction et de la ductilitr; inversement, dans les plaques 6paisses, l'augmentation de drlaminations dans l'rpaisseur (lires/t la faiblesse des proprirtrs tranverses) aux basses temprratures favorise localement des conditions de contrainte plane, renforgant par consrquent la tenacit6 par relaxation de la contrainte triaxiale. Dans les trles d'rpaisseur intermrdiaire cependant, l'absence de telles drlaminations dans l'rpaisseur permet l'apparition de la transition attendue des conditions de contrainte plane ~i celles de drformation plane; il en rrsulte que la trnacit6 diminue maintenant lorsque la temprrature drcroft.

Zusammenfassung--Der Zusammenhang zwischen Festigkeit und Z/ihigkeit wird bei den modernen Aluminium-Lithium-Legierungen 2090, 8090, 8091 und 2091 bei tiefen Temperaturen in Abh/ingigkeit von der Mikrostruktur, der Plattenorientierung und der Produktform (Platte oder Blech) untersucht. Besonderer Wert wird auf die Mikromechanismen des RiBfortschritts gelegt. Mit fallender Temperatur, von 298 zu 78 und 4 K, nehmen bei allen Legierungszusammensetzungen, Mikrostrukturen, Plattenorientierungen und Produktformen Festigkeit, Dehnung und Verfestigungsexponent zu. Allerdings kann die longitudinale (L-T, T-L) Z/ihigkeit ansteigen oder abfallen, je nach dem vorwiegenden mikroskopischen Mechanismus (Zusammenwachsen von Poren oder transgranulare Scherung) und der makroskopischen Mode (eben Dehnung oder ebene Spannung) des Bruches. Im allgemeinen ergeben Legierungsmikrostrukturen, bei denen sich bei teifer Temperatur Bruchmechanismus oder Bruchmode /indern, einen Abfall in der L-T-Z/ihigkeit. Dagegen h/ingt der beobachtete Zusammenhang zwischen Z~higkeit und Temperatur sehr stark vom Grad der sich an der Ril3spitze entwickelnden lokalen Spannungs-Dreiachsigkeit ab, wenn sich der Bruchmechanismus zwischen 298 und 4 K nicht 5.ndert. In sehr diinnen Blechen, in denen der Bruch-

2309

2310 VENKATESWARA RAO and RITCHIE: CRYOGENIC FRACTURE-TOUGHNESS OF A1-Li

mechanismus von ebener Spannung bestimmt ist, hfingt die Zunahme der Z~higkeit bei tiefer Temperatur mit einer einhergehenden Zunahme der Zugfestigkeit und der Duktilitfit zusammen. Andererseits f6rdert in dicken Platten das hS.ufigere Auftreten von Auftrennungen quer durch die Dicke bei tiefer Temperatur Iokal ebene Spannungsbedingungen, wobei die Zfihigkeit wegen der relaxierten dreiachsigen Bedingungen erh6ht wird. In Blechen mittlerer Dicke erm6glicht das Fehlen solcher Auftrennungen quer durch die Dicke den t2bergang yon ebenen Spannungs- zu ebenen Dehnungsbedingungen; im Ergebnis nimmt jetzt die ZS.higkeit mit fallender Temperatur ab.

1. INTRODUCTION

For the past 30 years, high-strength aluminum alloys have been the prime structural materials for many cryogenic applications because of their high specific strength, nonmagnetic behavior, stable microstructure, and retention of strength, ductility and toughness at very low temperatures [1]. For example, precipitation-hardened alloys containing up to 6.5 wt% Cu (2014-T651, 2021-T851, 2219-T851) or annealed solid-solutions with ~ 4.5 wt% Mg (5083-0) are widely used for storage and handling of many liquid cryogens including hydrogen, oxygen, nitrogen, helium and natural gas. In fact, the main rocket- propulsion fuel tanks on the Space Shuttle are fabricated from an heat-treated AI-Cu 2219-T87 alloy.

Recently, aluminum-lithium alloys have gained considerable attention in this regard, following reports [2-9] of remarkable improvements in their strength-toughness relationships at liquid-nitrogen (77 K) and liquid-helium (4 K) temperatures compared to room temperature (298 K). Because of this, aluminum-lithium alloys, which were originally developed as ultra-low-density materials to replace 2000 (AI-Cu-Mg) and 7000 (A1-Zn~2u-Mg) series aluminum alloys for aircraft structural use, are now additionally targeted for the fabrication of liquid-hydrogen and liquid-oxygen fuel tanks on future trans-atmospheric and hypersonic vehicles, including the National Aero-Space Plane [10].

While the low-temperature properties of aluminum-lithium alloys are clearly attractive, the underlying mechanisms that accurately describe their low-temperature toughness behavior, with respect to microstructure, specimen orientation and wrought product form, are not completely clear. Accordingly, the present investigation is aimed at obtaining an understanding of the principal micro-mechanisms governing cryogenic fracture toughness in these alloys.

2. BACKGROUND

Face-centered cubic (f.c.c.) materials, including AI and Cu alloys, austenitic stainless steels etc. are generally known to retain their ductility and toughness, with moderate strength improvements, at low temperatures. Moreover, since these materials rarely show evidence of cleavage fracture, even at cryogenic temperatures, they do not exhibit a clear ductile-brittle transition, unlike body-centered cubic (b.c.c.) structures. In fact, because transgranular microvoid coalescence is a common fracture mechan-

ism, even at 4K, the toughness and tensile- elongation properties of most f.c.c, metals tend to increase with strength, i.e. with reduction in temperature. These trends are illustrated for high-strength aluminum alloys in Fig. 1 [11-13], including recent results for aluminum-lithium alloy 2090-T81 [2, 5].

Although improved strength-toughness properties at cryogenic temperatures are not unique to AI-Li alloys, the magnitude of the increase, specifically in fracture toughness K~¢, is particularly striking. First results [2] on near peak-aged 2090-T81 (A1-2.1 Li-2.9Cu-0.1Zr, in wt%) plate showed nearly a 100% increase in fracture toughness (L-T orientation), combined with a ~45% increase in strength, a ~65% increase in tensile elongation and a ~ 12% increase in elastic modulus, between 298 and 4 K (Fig. 2). Subsequent studies [4-6, 9, 15], however, established that enhanced low-temperature toughness properties are not observed for all AI-Li alloys, nor for all specimen orientations, microstructural conditions and wrought product forms.

Results [2-9] to date have revealed that the elevation in fracture toughness of AI-Li alloys at cryogenic temperatures is principally observed in thick-section products, namely rolled plate. In contrast, thin (~ 1.6 mm) sheet of nominally similar composition and temper can exhibit a marked decrease in toughness at low temperatures, despite improvements in strength, ductility and strain-hardening rate, similar to plate material [15]. Moreover, the observed increase in low-temperature toughness, even for AI-Li alloy plate, is limited to the longitudinal and long- transverse (L-T, T-L, L + 45 °) orientations that lie in the rolling plane; short-transverse (S-L, S-T) orientations, conversely, show a small decrease [4-6]. Furthermore, improved cryogenic toughness properties are not common to all A1-Li alloys and microstructural conditions; in alloys where a transition in failure mechanism occurs at low temperatures, fracture toughness has been observed to decrease [6].

Several theories have been proposed for the superior cryogenic toughness of A1-Li alloys. As described in Refs [3-8], the phenomenon has been related to freezing of Na, K, and H-rich low melting- point impurities (or eutectic phases) at grain boundaries [3], to increased homogeneity of plastic deformation from higher strain-hardening rates [7] and reduced strain localization in wider and more closely spaced planar slip-bands [8], and to enhanced tensile ductility [16] at low temperatures, all mechanisms which fail to explain the concurrent reduction in S-T toughness of plate and reduction in L-T

VENKATESWARA RAO and RITCHIE: CRYOGENIC FRACTURE-TOUGHNESS OF A1-Li 2311

7 0

--.. 6 O

#.

5 O cj

40

~ 3o

2J 2o

o y ( k s i ) 5 0 6 0 7 0 8 0 9 0

I t I I

Conventional high-strength aluminum alloys (298 K)

~7 4K K

2090-T81

2 0 2 1 - T 8 ~ ~ " / / / /

298 K 7075-T651

tO0

i

6O

5O

4o

3O ~

2 0

10

7 0 0

10 I I I

300 400 500 6oo Tensile Yield Strength, (~ y (MPa)

Fig. 1. Strength-toughness relationships for commercial aluminum-lithium alloy 2090-T81 plate at 298, 77 and 4 K, compared to other high-strength aluminum plate alloys 2219-T87, 202 l-T81 and 7075-T651 [11-13]. Ambient temperature data on conventional high-strength aluminum alloys are shown for

comparison [14].

toughness of sheet. However, an alternative expla- nation [4--6], which accounts for toughness variations in all orientations and product forms, has focused on the greater tendency at low temperatures for delamination cracking (perpendicular to the rolling plane) along the weak, elongated grain structures, that are characteristic of wrought, thick-section forms of these materials (crack-divider delamination toughening).

Although most of these mechanisms may act in concert to varying degree, delamination toughening appears to provide a primary basis for many of the observations reported on the cryogenic-toughness behavior of AI-Li alloys. Akin to behavior in laminated composites, delamination-type separation along the weak and elongated, short-transverse grain boundaries can actually enhance toughness in the crack-divider and crack-arrester orientations (Fig. 3). Specifically, toughening in A1-Li alloy plates can be achieved at low temperatures in the L-T, T-L and L + 45 ° (crack-divider) orientations by the relaxation in stress triaxiality and resulting transformation of the global plane-strain fracture into a series of local plane-stress ("thin-sheet") failures, due to the formation of through-thickness delamination cracks ahead of the crack tip [4-6]; additionally, the elevation in low-temperature strength, work- hardening rate and ductility can further accentuate this toughening effect. Accordingly, in the absence of such intergranular spitting, e.g. in thin sheet where through-thickness stresses are negligible, increase in toughness at low temperatures may not be seen [15].

The occurrence of such short-transverse splitting in A1-Li alloys can also promote toughening in the T-S and L-S (crack-arrester) orientations by blunting and deflecting the crack (typically through ~90 ° ) along the weak, grain-boundary interfaces, i.e. the Cook-Gordon mechanism [17] for crack stopping at weak interfaces (Fig. 3). This delamination in A1-Li alloys is consistent with the T-S toughness being at least four times the (S-T) toughness of the interface (see results for 2090-T81 in Ref. [6]), in accordance with proposed criteria for crack deflection at interfaces [18]. Moreover, the measured T-S toughness in 2090-T81 plates is 65 MPa~/m [6], approximately twice the L-T toughness, again consistent with a reduction in the local stress intensity at the deflected crack tip by a factor of ~ 2 due to an in-plane tilt of the crack by ~90 °, predicted using simple crack- deflection mechanics [6, 19].

Toughening by the introduction of aligned, weak interfaces is not uncommon in other materials; it has been utilized for a number of years in the design of laminated steels [20-24], graphite-epoxy laminates [25] and advanced ceramic composites [26]. More recently, these mechanisms have been associated with the fracture behavior of solid-state- welded Damascus steels [29], high-temperature A1-Fe alloys containing Ce, Mo and V [28], and even sea shells [29]. In virtually all these examples, the inclusion of planes of weakness in the microstructure, either by thermomechanical processing in monolithic alloys or hybridization in composites,

A M 38/11--S


900

800

0 0.

700

o 600

500

400

°C -200 -100 0

( a ) I '2090- T81 12.7 mm thick plote

" ~ ° _

Ductility

i t I I

g 2

2O tlJ

15

10

5

70~ ( b ) l I I J

60

~ 20 .~

10 --~ S-L, s-r

u I o lOO 200 300

Temperoture, K

Fig. 2. Variation with temperature of (a) uniaxial yield strength, tensile strength, tensile elongation in the longitudinal (L) orientation, and (b) plane-strain fracture-toughness in the L-T, T-L, L+45 °, S-T, S-L orientations for 12.7mm thick AI-Li-Cu-Zr alloy 2090-T81 plate. Data

taken from Refs [2, 5].

has resulted in Substantial toughening along specific orientations.

As these explanations are still the subject of some controversy for A1-Li alloys, the objective of the present paper is to critically reexamine the role of crack-divider delamination toughening in influence- ing the low-temperature mechanical behavior of these materials. The approach is to compare the fracture toughness (and uniaxial tensile properties) of commercially-rolled plate with corresponding behavior in commercial and artificially-machined thin sheet (where through-thickness splitting is not expected to occur), as a function of test temperature, microstructure and specimen orientation, in order to elucidate the primary mechanisms of cryogenic toughening.

3. EXPERIMENTAL

3.1. Materials and microstructures

Near peak-aged A1-Li-Cu-Zr alloy 2090 was examined in the form of commercially-rolled 12.7 mm thick plate (T81) and 1.6 mm thin sheet (T83). The T83 and T81 temper designations refer to thermomechanical treatments, which typically involve solution-treating at 549°C, water quenching, and a 6% permanent stretch prior to artificially aging for 24 h at 163°C. In addition, 11-16-mm thick plates of A1-Li -Cu-Mg-Zr alloys 8090, 8091 and 2091 were tested in the underaged (T351) and peak-aged (T8) tempers. Chemical compositions are listed in Table 1;

-S

rester

L-T z

Crack di'

Fig. 3. Three-dimensional micrograph of the grain structure of 2090-T81 plate material, showing the orientation of test samples in the L-T (crack divider), T-S (crack arrester) and S-T (crack delamination)

orientations.

VENKATESWARA RAO and RITCHIE: CRYOGENIC FRACTURE-TOUGHNESS OF A1-Li

Table 1. Nominal chemical compositions of aluminum-lithium alloys investigated (wt%)

2313

Alloy Li Cu M g Zn Fe Si Ti Zr AI

2090 2.05 2.86 0.01 0.01 0.02 0.01 0.02 0.12 bal. 8090 2.50 1.30 0.70 - - 0.20 0.10 - - 0.12 bal. 8091 2.60 1.90 0.90 - - 0.10 0.10 - - 0.12 bal. 2091 1.7-2.3 1.8-2.5 1.1-1.9 0.25 0.30 0.20 0.10 0.044).16 bal.

thermomechanical treatments used to obtain the different microstructures are detailed elsewhere [30].

In the plate materials, grain structures for all alloys were unrecrystallized with large pan-cake shaped grains, roughly 25-50 pm thick, 65-500 ~tm wide and elongated several millimeters in the rolling direction (e.g. Fig. 3); some degree of recrystallization, however, was apparent in 2091. Corresponding aspect ratios, computed as the ratio of grain length or width to thickness, are between 10-20 and 5-10, in the longitudinal and transverse directions, respectively. In contrast, fine equiaxed subgrain structures were seen in 2090-T83 sheet, indicative of continuous recrystallization within the elongated grains. Owing to prior cold-working, plate and sheet products of all alloys exhibit strong deformation textures; these textures have been found to be predominantly of the brass-type ({123}(112)), with evidence of weaker S ({123}(634)) and copper ({112}(111)) types [31-33]. Accordingly, mechanical properties are strongly an- isotropic and vary with specimen orientation, in and out-of the rolling plane and through the plate thickness [30, 34].

Microscopically, hardening in peak-aged 2090 was achieved by homogeneous matrix distributions of spherical, coherent, ordered 6' (A13Li) precipitates, 0'-like (A12Cu) and T l (AI2CuLi) plates, with fl' (A13Zr) dispersoids. Limited heterogeneous precipitation along grain and subgrain boundaries was also evident, resulting in 6'-precipitate-free-zones (PFZs), which were ~ 50 and 100 nm wide in sheet and plate, respectively. Similar features were noted for Mg- containing 8090 and 8091 alloys in the T8 temper, except that 0 ' and T1 plates in the matrix were replaced by S (A12CuMg) laths. Small (~500nm) 6'-PFZs were present in 8090 as a consequence of grain-boundary precipitation; heterogeneous precipitation was more extensive in 8091-T8 and PFZs were correspondingly wider (~ l#m). Conversely, natu- rally-aged T35t microstructures of all alloys were hardened by solid-solution strengthening and very fine distributions of 6' spheres and fl' dispersoids, similar to peak-aged 2091. In addition, micron-size undissolved intermetallic or constituent particles, re- tained during homogenization and rich in Fe, Cu and Mg, were distributed as stringers along high-angle grain boundaries in all alloys.

3.2. Mechanical testing

Fracture-toughness tests on plate material were conducted using 6-16mm thick, compact tension [C(T)] and four-point, single-edged-notch bend [SEN(B)] samples for orientations in the rolling plane

(L-T, T-L, L-S, T-S, L + 45°), and with 6 mm thick double-cantilever beam (DCB) samples for the short- transverse (S-L, S-T) properties. Specimens were initially fatigue precracked, at a load ratio of 0.1 with the maximum stress intensity (Kmax) not exceeding 10MPax/m, to a crack length of roughly half the specimen width. Tests were performed at 298 K (ambient air), 196K (dry ice in ethanol) and 77 K (liquid-nitrogen) temperatures, using electro- servohydraulic testing machines operating under displacement control.

Corresponding fracture-toughness behavior in 1.6 mm 2090-T83 sheet was assessed at maximum load (plateau) by determining KR(Aa ) resistance- curves (R-curves), as a function of temperature, using fatigue precracked, compact C(T) specimens, machined in the L-T orientation. Crack initiation and growth was monitored using d.c. electrical poten- tial and elastic unloading compliance methods; suit- able buckling constraints were employed to minimize out-of-plane shear deformation. Additional tests were performed for thicknesses of 0.8, 0.34 and 0.24 mm, obtained by machining away the surface layers of fatigue-precracked, C(T) samples of 2090-T83 sheet.

Uniaxial stress-strain properties for plate material were evaluated in the longitudinal (L) direction using 6.4 mm round tensile specimens with 25 mm gauge length. Corresponding tests for sheet material employed 6.4 mm wide, fiat tensile specimens with identical gauge length. Ambient temperature mechanical properties for all alloys are included in Table 2.

3.3. Fractography

Macroscopic fracture modes and crack-path mor- phologies were examined on metallographic sections of broken and unbroken test pieces, taken through the specimen thickness, perpendicular and parallel to the crack-growth direction. Scanning electron microscopy was used to identify the microscopic fracture mechanisms. A distinction is made here between the macroscopic mode and microscopic mechanism of fracture, roughly based on the magnification level at which the observations are made. The macroscopic mode refers to whether failure (as observed on profiles taken across the specimen thickness) occurs under nominally plane strain ("flat" fracture), plane stress (shear lips or "slant" fracture), or a combi- nation of both (mixed mode or "slant plus flat" fracture). On the other hand, the microscopic mechanism relates to the local failure process (as identified by features on the separated surfaces using electron microscopy), e.g. microvoid coalescence or intergranular cracking.

2314 VENKATESWARA RAO and RITCHIE: CRYOGENIC FRACTURE-TOUGHNESS OF AI-Li

Table 2. Fracture toughness and tensile properties" of advanced aluminum-lithium alloys at 298 and 77 K

Fracture Plastic Yield strength U.T.S. % elongation toughness Strain hardening zone size

(MPa) (MPa) (on 25 mm) Kit (MPa~/m) exponent (n) ry (mm)

Alloy 298 K 77 K 298 K 77 K 298 K 77 K 298 K 77 K 298 K 77 K 298 K 77 K

11-16mm thick plate 2090-T81 552 587 589 642 11 14 36 5P (L-T) 0.06 0.15 0.7 1.2

17 15 (S-L) 65 b -- (T~S)

2091-T351 369 442 451 596 10 16 33 b 41 b (L-T) 0.12 0.22 1.3 1.4 19 17 (S-L)

2091-T8 425 483 481 610 8 14 46 b 44 b (L-T) 0.10 0.21 1.9 1.3 25 16 (S-L)

8090-T351 226 256 352 486 17 24 27 b 20 (L-T) 0.19 0.34 2.3 1.0 16 17 (S-L)

8090-T8 482 517 534 638 6 8 36 38 (L-T) 0.80 - - 0.9 0.8 13 12 (S-L)

8091-T351 309 382 417 572 II 14 38 b 28 (L-T) 0.16 0.28 2.4 0.9 17 10 (S-L)

8091-T8 537 574 581 697 6 12 20 38 (L-T) 0.07 0.18 0.2 0.7 9 8 (S-L)

2090-7"83 sheet 1.6 mm 505 568 549 674 6.8 8.0 43 ¢ 30 ¢ (L-T) 0.05 0.08 0.8 0.4 0.75 mm 34 c 35 ¢ (L-T) 0.7 0.6 0.34 mm 30 ¢ 40 c (L-T) 0.5 0.8 0.24 mm 22 ¢ 42 ¢ (L-T) 0.3 0.9

aL direction. bK¢ values not meeting the ASTM plane-strain thickness criterion. eR-curve plateau "plane-stress" fracture toughness K c values.

The extent o f through-thickness delamination was evaluated fractographically and by metallographic sectioning ahead of the unbroken notch in double- edge notched samples, tested under four-point bending, both at 298 and 77 K. Sectioning was performed perpendicular to the crack front in increments of 100/~m; at each section, the height (or depth) and spacing of individual delaminations were recorded. In addit ion, fracture-toughness samples were sectioned at loads corresponding to 50, 85 and 90% of Klc, to discern whether such delaminations occur ahead of the crack tip, prior to or following, crack initiation.

4. RESULTS

4.1. Tensile properties

Uniaxial tensile properties of the plate 2090, 2091, 8090 and 8091 alloys and 2090 sheet material are shown in Figs 4(a)-6(a) as a function of temperature; additional data are listed in Table 2. The yield and tensile strengths, elongation and strain-hardening exponent increase with decreasing temperature for all alloys, independent of composi t ion and microstructure, al though the effect is somewhat larger in plate material. Compared to room temperature, yield and tensile strengths are ~ 10 to 40% higher and elongation values up to 75% higher at 77 K. Low- temperature fractures were seen to occur at maximum load with little evidence of localized necking before final fracture.

4.2. Fracture-toughness properties

Corresponding plane-strain fracture-toughness data for plate materials and (plateau) toughness from R-curves on sheet material are plotted as a function of test temperature in Figs 4(b)-6(b), with additional

data listed in Table 2. It is clearly apparent that the toughness of thick plates is strongly dependent upon specimen orientation. At both ambient and cryogenic temperatures, K~¢ values in plate material are

*C -200 -tO0 0

7 0 0 | (a) i l ,

I 600 r

5 0 0 Jstrength o ~ ~'~ I YieLd

[Ductil ity ~__ - 25 UJ

~ ~ . . . . . - 20

- 1 5

1oo I ""~--~-~-~---~. 1o ol I I ~ I 5

"E 70 --> o ( b ) • 2o91 -T8 n 6 0 A 8 0 9 0 - T 3 5 1 =} [] 8091 - T 3 5 1

50 s- --s

. 40 {

3 0

I I I o 0 100 200 300 LL

Temper0ture, K

Fig. 4. Temperature dependence of (a) yield strength and tensile elongation in the L orientation, and (b) fracture toughness in l l -16mm thick 8090-T351, 8091-T351 and 209 l-T8 plate in the L-T and S-L orientations. Note that for these alloys, the fracture toughness decreases with temperature despite increases in strength and elongation.

VENKATESWARA RAO and RITCHIE: CRYOGENIC FRACTURE-TOUGHNESS OF AI-Li 2315

*C - 2 0 0 - 1 0 0 0

8 0 0 1 ( 0 ) I I I /

7 0 0

6 0 0 F - - - - - - : L - - - ~ _ _ _

i 400~ "~ I o 300 L- - 20 ~

200 I ~ -~ ~ - 15

lOO ~- %_ ~ _ : ~ = ~ Z _ _ - ~ - , o

- r o L ( b ) • 2 0 9 0 - 1 8 1 "~P F ~ ~. 2091 -T351 m 60 o. I ~ [] 8090 - T 8 '5 50 I ;'~e o 8091 - T 8

40 L-T ~ 30 ~ 2 0

10 L

0 100 200 300

" Temperature, K

Fig. 5. Temperature dependence of (a) yield strength, tensile strength and tensile elongation in the L orientation, and (b) fracture-toughness behavior in 11-16 mm thick 2090-T81, 8090-T8, 8091-T8 and 2091-T351 plate in the L-T and S-L orientations. Note that for these alloys, the strength, tensile ductility and L-T fracture toughness all increase with

decreasing temperature.

~20-80% lower in the short-transverse (S-L, S-T) orientations compared to the more commonly tested long-transverse (L-T or crack-divider) orientation. Also as noted above, the toughness of 2090-T81 plate, at 298 K in the T-S (crack-arrester) orientation, is 65 MPa~/m, over 80% higher than in the L-T orientation.

The temperature-dependence of these properties is also a function of orientation (Table 2). Short- transverse toughness for all plate alloys shows a small decrease at low temperatures; fracture surfaces involve delamination or lameUar-type splitting (e.g. as in wood) along grain boundaries, with highly linear crack paths. In contrast, the L-T fracture toughness can increase or decrease with reduction in temperature, depending upon the prevailing fracture mechanism and mode of failure; behavior in plate and sheet is described below.

4.2. I. Behavior in thick plate. As noted in Ref. [6], microstructures which exhibit a transition in fracture mechanism with decrease in temperature (Fig. 7), i.e. peak-aged 2091 and underaged 8090 and 8091, display lower (L-T) toughness at cryogenic temperatures, despite elevated strength, ductility and strain hardening (Fig. 4). Fractographically, mixed-mode ("flat" with large shear lips) fracture surfaces at 298 K [Fig. 7(a)], involving transgranular microvoid co-

*C - 2 0 0 -100 0

8OO I I I ( a ) 2 0 9 0 - r e 3

1.6 mm thin sheets

I~. 7 0 0 c

6 0 0 ..n._Yiel d s t r e n g t h e n _ - 10

- ~ ~e,. ILl

5 0 O

4 0 0 I I I 5 ~E 60

, ~ r i i ~l.sf

:E

50 I"

:~ 40--

30-- I - I

20 I I I 0 1 O0 200 3 O0

T e m p e r a t u r e , K

Fig. 6. Temperature dependence of (a) yield strength and tensile elongation in the L orientation, and (b) L-T fracture toughness in commercial 1.6 mm thin 2090-T83 sheet. Note that unlike the 2090-T81 plate, the toughness decreases with decreasing temperature, despite increases in the strength and

tensile ductility.

alescence around cracked or uncracked Fe- and Cu-rich intermetallics and constituent particles [Fig. 7(c)] at room temperature, become fully "flat" (plane strain)' fractures at 77K [Fig. 7(b)] and involve a brittle transgranular-shear type separation, with occasional delamination cracks perpendicular to the fracture plane [Fig. 7(d)].

Conversely, microstructures which undergo no such fracture-mechanism transition with decrease in temperature (Fig. 8), i.e. peak-aged 2090, 8090 and 8091 and underaged 2091, show remarkably higher (L-T) toughness at cryogenic temperatures [Fig. 5(b)]. At both 298 and 77 K, these alloys show a coarse transgranular shear fracture mechanism, with limited regions of microvoid coalescence interdispersed between intergranular delamination cracks that run parallel to the crack-growth direction (Fig. 8). Although the basic fracture mechanism is unchanged with temperature, the incidence and depth of through-thickness delamination cracks is significantly greater at 77 K [Fig. 8(c, d)]. Corresponding fracture paths, as observed along the direction of crack growth, are predominantly transgranular and do not exhibit significant crack branching, deflection or bifurcation at the onset of instability (K = K~c), unlike behavior in experimental AI-Li -Cu alloys [35]; such deviations are largely confined to the free surfaces and not seen in the specimen interior [30].


Fig. 7. Micrographs illustrating (a, b) macroscopic fracture modes and (c, d) fracture mechanisms at 298 and 77 K, typical of 8090-T351, 8091-T351 and 2091-T8 alloys that exhibit reduced toughness at lower temperatures. Micrographs obtained from 11 mm thick 2091-T8 plate in the L-T orientation and indicate a fracture-mechanism transition from microvoid coalescence to transgranular shear with decrease in

temperature. Arrow indicates the general direction of crack advance on fracture surfaces.

4.2.2. Behavior in thin sheet. In contrast to the L - T fracture toughness of the 2090-T81 (unrecrystallized) plate, which increases sharply at cryogenic temperatures (Fig. 5b), corresponding behavior in commercial 1.6 mm thin (partially recrystaUized) 2090-T83 sheet, of nominally identical composit ion and

microstructure, shows quite different behavior; in fact the sheet material displays decreased toughness at cryogenic temperatures, despite similar improvements in strength, elongation and strain-hardening exponent (Fig. 6). At 298 K, toughness increases with crack extension (rising R-curve behavior) as the

298 K

77 K

Fig. 8. Micrographs illustrating (a, b) macroscopic fracture modes and (c, d) fracture mechanisms at 298 and 77 K, typical of 2090-T81, 8090-T8, 8091-T8 and 2091-T351 alloys that exhibit increased L T toughness at lower temperatures. Micrographs obtained from 12.7 mm thick 2090-T81 plate and indicate enhanced through-thickness (short-transverse) splitting with decrease in temperature. Arrow indicates the

general direction of crack advance on fracture surfaces.


spread of plasticity relaxes through-thickness constraint to give a fully plane-stress (~45 ° "slant") fracture [Fig. 9(a). Failures at 77 K, on the other hand, appear less ductile and are mixed mode ("slant plus flat"), approaching plane-strain conditions due to the elevation in strength [Fig. 9(b)], However, an intergranular-ductile microvoid coalescence fracture mechanism prevails at both 77 and 298 K [Fig. 9(c, d)], with no evidence of through- thickness delamination cracking.

The variation in sheet toughness with temperature, however, is a strong function of specimen thickness. Measured toughness values in artificially-machined 2090-T83 sheet, for a range of thicknesses from 1.6 to 0.24 mm, are plotted as a function of temperature in Fig. 10; results for 12.7 mm thick 2090-T81 plate are shown for comparison. Corresponding changes in fracture mode are documented in Fig. 11. It can be seen that for completely unconstrained failures where the fracture mode remains in "plane stress" ("slant"

fracture) between 298 and 77 K, i,e., for sheet thicknesses less than 0.34mm, toughness increases with decreasing temperature. Thus, although the strength, tensile ductility and strain hardening all increase with decrease in temperature, the variation in toughness is dependent on the macroscopic fracture mode, i.e. the degree of local stress triaxiality that develops at the crack tip; the mechanism of fracture, however, remains microvoid coalescence throughout.

5. DISCUSSION

5.1. Mechanisms for improved cryogenic toughness

The current results, and those published previously [2-9, 15], provide clear evidence that the marked improvement in fracture toughness at cryogenic temperatures is not a general characteristic of all AI-Li alloy microstructures, nor of all specimen orientations or wrought-product forms. First, plate alloys that undergo a fracture-mechanism transition (8090-

a 200 om I I

298 K

77 K

Fig. 9. Micrographs illustrating (a, b) macroscopic fracture modes and (c, d) fracture mechanisms at 298 and 77 K, in 1.6 mm thin 2090-T83 sheet (L-T orientation). Note the fracture-mode change from a 45 ° slant fracture at 298K to mixed mode failure at 77K, although the fracture mechanism remains

unchanged. Direction of crack advance on fracture surfaces is denoted by the vertical arrow in (c).


=~ 4o

3 0

- 200 70 I

a 12.7 m m thick 60 . " " , , 2 0 9 0 - T 8 1 Plate

%%%/

20

o C

-I00 0

2 0 9 0 - T 8 3 Sheet o B = 1.60 m m

• B - - 0 .76 m m

B - - 0 . 3 4 m m

• B = 0 .24 m m

1 0 I I I

0 100 2 0 0 3 0 0 Temperature, K

Fig. 10. Comparison of the temperature dependence of L-T fracture toughness for 12.7 mm thick plates and 0.2-1.6 mm

thin sheets of near peak-aged 2090-T8 alloy.

T351, 8091-T351, 2091-T8) do not exhibit enhanced L-T toughness at lower temperatures. Second, it is only in the plate alloys (2090-T81, 8090-T8, 8091-T8, 2091-T3) where the fracture mechanism (transgranular shear) remains unchanged that L-T plane-strain fracture toughness values show a significant enhance- ment at 77 and 4 K compared to 298 K. Third, the corresponding "non plane-strain" fracture toughness of 2090-T83 sheet material can show either a decrease or increase with temperature, depending upon the local triaxial constraint that develops at the crack tip at various temperatures, i.e. the temperature dependence of toughness is a function of thickness, even though the fracture mechanism (microvoid coalescence) remains unchanged. Finally, the short- transverse (S-L, S-T) toughness of all plate alloys shows a small decrease at low temperatures [4--6].

Such results are inconsistent with many published explanations for the cryogenic toughness behavior of AI-Li alloys, which relate the low-temperature increase in L-T toughness of plate alloys solely to the solidification of grain-boundary liquid phases [3] or with increased slip homogeneity and the resulting increase in strain-hardening rates and tensile ductility [7, 8]. While these mechanisms may well provide a contribution to the enhanced toughness, they fail to explain why 2090-T81 plate and T83 sheet material, which have nominally identical chemical composition, similar thermomechanical treatments and strengthening precipitates, and display increased strength, elongation and strain hardening at lower temperatures, show opposite trends in toughness behavior with decrease in temperature (in each case with no change in fracture mechanism). Moreover, these theories do not explain why the toughening is specific to specimen orientations in the rolling plane

(L-T, T-L, L+45°) , and not to S-L and S-T orientations. In fact, increased toughness from the freezing of grain-boundary liquid phases at low temperatures would be expected to be most pronounced in the latter orientations, where failures are totally intergranular.

The arguments for improved L-T toughness at cryogenic temperatures that rely on increased ductility and strain hardening [7, 8, 16], are largely based on expressions derived by various micromechanical models for critical-strain controlled fracture [36--42]. Essentially, all these models are variations of the strain-across-distance formulation as outlined in the Appendix, and imply that toughness is proportional to strength, ductility and a characteristic microstructural dimension; strain hardening is incorporated in a rather empirical fashion [36, 38]. Although, these models provide some justification for the elevation in Kit at low temperatures, due to increases in strength and tensile elongation, quantitatively they have had limited success [16, 30]. This may be due to one of several reasons: (i) the models are dependent on local fracture parameters i.e. fracture strain measured under a highly triaxial stress-state (see Appendix), (ii) they do not take into account the change in constraint at the crack tip due to splitting, and (iii) there is no definitive proof that mechanisms of L-T fracture in advanced AI-Li alloys are, in fact, strain-controlled.

Conversely, explanations which include the effects of crack-divider delamination toughening [4--6] in plate material appear to be far more convincing. Apart from plate alloys where the reduction in toughness at cryogenic temperatures is associated with a distinct transition in the fracture mechanism from ductile microvoid coalescence to transgranular shear (Figs 4 and 7), the elevation in L-T toughness observed in plate alloys that show no such transition can be attributed to the greater propensity for through-thickness (short-transverse) splitting, normal to the fracture plane, at low temperatures (Figs 5 and 8). Since the short-transverse splitting occurs by a brittle grain-boundary decohesion fracture [6-9], one can reasonably model the local failure process as being stress-controlled [41,43]. This mechanism would thus be promoted in thick (plane-strain) sections by the development of a finite through-thickness stress (tr=), and at low temperatures by the increase in yield strength (try) and strain-hardening exponent (0 < n < 1), both of which elevate the magnitude of the local tensile stresses developed at the crack tip. The latter fact is consistent with the observed decrease in measured short-transverse toughness at low temperatures (Fig. 2). In addition, tryz and tr= shear stresses, that act to slide the grains past one another, may also contribute to splitting due to the inherently poor shear resistance of laminated microstructures.

Micro-mechanical models for such stress-controlled fracture (see Appendix) have been well developed for brittle transgranular (cleavage) fracture [41, 43] and somewhat less so for intergranular crack-


Thickness = 1.6 mm 0.75 0.34 0.24

2319

2 9 8 K

2 0 0 K

77 K

Fig. 11. Optical micrographs showing the variation in macroscopic fracture mode (L-T orientation) with test temperature and specimen thickness in 2090-T83 sheet. Crack-growth direction is given by the

outward normal to the plane of the paper.

ing [44, 45]. In either case, the toughness is predicted to be of the form [41]

gic~ [{~r~ < }(1 +n)12n)/{tTy }(1 - n)/En]{l~ }1/2 (1)

where n takes values between 0 and 1, tT* is the local fracture stress, and l* is the characteristic microstructural dimension (typically the average

spacing between crack-initiating constituent phases or precipitates, in this case distributed along high- angle grain boundaries [i.e. of order of the grain thickness along the short-transverse direction)]. Lacking data for the temperature variation in yield strength, fracture stress and work-hardening exponent in the short-transverse orientations, equation (1)


cannot be evaluated quantitatively; however, concep- tually the model is consistent with the observed decrease in S--T toughness and increased tendency for splitting in plate material at low temperatures. We now examine the role of such splitting on the L-T fracture-toughness behavior.

5.2. Role of crack-divider delamination toughening

5.2.1. Behavior in thick plate. In the absence of through-thickness splitting, fractures in 12.7mm plate material would be expected to be under plane- strain conditions at both 298 and 77 K as plastic- zone sizes (estimated from ry "~ { 1/2~ } {Kit/ay }2) are some 10~0 times smaller than the plate thickness. Such highly triaxiai conditions elevate the crack-tip tensile stresses to between three to five times the yield stress (depending upon the yield strain and strain- hardening exponent), and can result in the local fracture strain (ductility) at the crack tip to be typically 3 to 10 times smaller than the uniaxial tensile ductility (see Appendix). Through-thickness stresses under fully plane-strain conditions may be estimated to be roughly 2-3 times ay, based on the Tresca criterion for yielding (~ryy- axx = ay) and the incompressible (zero volume change) nature of plastic deformation [a= = 0.5(ax~ + O'yy)]. The principal effect of the increased splitting at cryogenic temperatures is to relieve the through-thickness stresses (a~z ~ 0) and thus to reduce this constraint at the crack tip. Accordingly, deformation conditions within the plastic zone are transformed from plane strain (fully constrained) to a "natural laminate" of near plane stress ligaments (unconstrained) at low temperatures,

with a corresponding reduction in the magnitude of crack-tip tensile stresses and an increase in the required local ductility for fracture. Consequently, the nominally-flat mode of fracture seen at 298 K is replaced by a series of individually-slant fractures at 77 K [Fig. 8(a, b)]. In fact at 77 K, the spacing and height of the major delaminations (Figs 12 and 13) is comparable to the plane-stress ligament size, i.e. of order of the plastic-zone size. Since it is well known that the plane-stress fracture toughness K, is higher than the plane-strain Kxc value (room-temperature measurements [6] give values of 45 and 36 MPax/m, respectively, in 2090-T81 plate), the through- thickness splitting phenomenon would be expected to yield increased toughness.

A lower-bound estimate [6] for the magnitude of this toughening effect (ignoring the energy required for delamination itself and improved strength at low temperatures) can be derived by assuming that the material behaves like a laminate, where each of m ligaments sees a constant tensile stress. At failure, each ligament carries a load P, given in terms of the full specimen width and thickness B, by

P = (K'~ B W m)/[mf(a / W)] (2)

where f ( a / W ) is a function of the crack length-to- width ratio for the specific specimen geometry and K~ is the fracture toughness of a ligament of thickness B/m. Since the total load carried by the full thickness is mP, the fracture toughness of the laminate will be K ' . Since the major delamination spacings appear to be of the order of the plane-stress ligament size, K c

¢

Fig. 12. (a) Schematic illustration of crack-divider delamination toughening via through-thickness splitting, and sections taken through the specimen thickness ahead of the fatigUe-precrack tip, (b) at a load corresponding to 0.9 Po, where PQ is load at crack initiation and (c) following fast fracture. Optical

micrographs obtained from near peak-aged 2090-T81 plate tested at 77 K.


ROC " (29~

2321

~ i q t . . . . . .

err

T-Y

Fig. 13. Three dimensional maps of the height (measured as length above the main fracture plane) and spacing of (through-thickness) delaminations ahead of the crack tip at (a) 298 K and (b) 77 K. Data obtained by serially sectioning the unbroken ligaments of fatigue-precracked, double-notched test specimens of 2090-T81 plate (L-T) loaded in four-point bending. Only regions above the main fracture

plane are plotted for clarity.

should be of the order of the plane-stress fracture toughness K~.

It is important to note that for delamination to be effective in increasing the crack-initiation fracture toughness, Klc , the through-thickness splits must be formed prior to initiation of the main crack. To verify

this hypothesis, a series of L-T fracture-toughness specimens were monotonically loaded at 77 K to various loads before the onset of quasi-static crack growth and metallographically sectioned perpendicular to the main-crack plane in the region directly ahead of the precrack. As shown in Fig. 12, intergran-


x = 0.38 ram

Fig. 14. Optical micrographs of metallographic sections taken at distance of x = 0 to 381/~m from the crack tip in 12.7mm thick 2090-T81 plate at 77K. Note the presence of a large number of minor delaminations between the major splits, which progressively diminish with distance away from the crack

tip.

ular delaminations begin to form ahead of the precrack tip at roughly 85-90% of the crack-initiation load PQ. Subsequent crack advance following initiation is essentially catastrophic; little or no stable extension is seen, to the effect that the influence of splitting on crack-growth toughness is relatively minimal.

Another factor of importance is that some degree of splitting is often observed even in room- temperature fractures (Fig. 8). To examine this, the separation and length of the splits were measured throughout the specimen thickness for fracture- toughness tests on 2090-T81 plate at 298 and 77 K, from a series of optical micrographs taken at various distances ahead of the crack tip just prior to final failure using precracked, double-notched, four-point bend specimenst (Figs 14 and 15). It is apparent that a single delamination spacing and height is difficult to define since a distribution of delaminations exists with several smaller splits in between the major ones. Moreover, the number of delaminations present is a

tSince the bending moment is constant between the inner loading points in four-point bend, both precracked notches see nominally identical stress and deformation fields. At failure, sections taken ahead of the unbroken precracked notch thus reflect conditions just prior to instability.

function of position, with the maximum degree of splitting occurring some distance ahead of the crack tip where stress triaxiality is the largest. However, for the delaminations to be effective in reducing through- thickness constraint by permitting plane-stress 45 ° shear through each ligament (Fig. 15), the length (l) of the splits must be at least equal to their separation (s). At room temperature, this condit ion is not always satisfied as the average spacing beween delaminations is roughly I mm compared to a length of typically 0.6 mm (Fig. 13). Deformation at 298 K is thus reasoned to remain predominantly in plane strain, as the delaminations are of insufficient size to relax constraint effectively. Conversely, the spacing

t/// / /

1/ _ _ _ ~ / _ _ _ _

short-transverse Y delaminations /,_/~ /... //__/_/ _~ rm_aaitnu r.e#l_ ane

,a > s to relieve constraint F i g . 15. S c h e m a t i c m o d e l o f ~ 4 5 ° s h e a r t h r o u g h a d e l a m i - nated ligament, illustrating that the length (1) of the through-thickness splits must be at least equal to the split

spacing (s) for constraint to be relaxed.

VENKATESWARA RAO and RITCHIE: CRYOGENIC FRACTURE-TOUGHNESS OF AI-Li 2323

between the major splits at 77 K is considerably 'smaller (~ 0.3 mm), compared to an average length of split of the order of 1 mm (over the majority of the thickness). In this case through-thickness splitting will be effective in relaxing constraint such that crack-divider delamination toughening can occur.

The difficulty noted above of defining the spacing of major splits, which are effective in relaxing constraint, is also probably responsible for the apparent lack of a one-to-one correspondence between delamination spacing (measured on the fracture surface) and the fracture toughness, reported by certain authors [16]. In addition, theories for pure plane- stress slant fracture, such as those proposed by Cottrell and others [46, 47], predict that the toughness for such shear failures scales with (EayB) 1/2, where E is the elastic modulus and B is the thickness of the ligament (see Appendix). Thus, the effectiveness of through-thickness delaminations in enhancing toughness at cryogenic temperatures is a function of both the elevation in strength and stiffness with decrease in temperature, as well as the extent of intergranular splitting; a monotonic dependency between spacing and toughness is therefore unlikely.

5.2.2. Behavior in thin sheet. In contrast to behavior in plate material, no through-thickness short- transverse splitting was observed in 2090-T83 alloy sheet, as shown in Figs 9 and 11. This can be attributed to the lack of through-thickness stresses in the thin sections, and to the more equiaxed (continu- ously recrystallized) nature of the grain structure in the sheet material. For the commercial 1.6-mm thin sheet, a "plane-stress" (45 ° slant) fracture mode at 298 K becomes a mixed (plane-stress plus plane- strain) mode of fracture at 77K (Fig. 11), even though the fracture mechanism (microvoid coalescence) remains unchanged (Fig. 9). Due to diminished plastic flow at low temperatures, the crack-tip plasticity at 77 K is insufficient to completely relieve constraint; a part fiat, part slant fracture mode thus results, with a consequent decrease in toughness compared to room temperature. Similar to plate material, the plastic-zone size at fracture at both 298 and 77 K is comparable to one to two times the width of the slant fracture section, consistent with plane- stress conditions.

With progressively thinner sheet, however, the fracture-mode transition is suppressed such that conditions approach "plane stress" at both 298 and 77 K (Fig. 11). This is demonstrated by the behavior of mechanically-thinned 2090-T83 sheet where for thicknesses below 0.34 mm, the fracture mode remains 45 ° slant at cryogenic temperatures with the result that the toughness once again increases with decreasing temperature (Fig. 10), presumably due to the enhanced strength, tensile ductility and strain hardening of the alloy at 77 K. Note, however, that the room- temperature toughness decreases with thickness, consistent with Cottrell's model [46, 47] for plane-stress fracture where K c is proportional to (EayB) 1/2. This

can be seen in Fig. 16 by replotting the sheet toughness data as a function of thickness B (for various test temperatures and consequently different yield strengths); it is clear that, for a given temperature, there is a maximum in toughness at the largest thickness resulting in a fully-slant fracture at that temperature. Above this thickness, fully plane-stress conditions no longer prevail at the crack tip (except at the surface), and the toughness decreases with increasing constraint until it approaches the plane- strain value; below this thickness, deformation conditions remain totally unconstrained and the toughness decreases with B ~/2. With respect to the thick-plate material, this implies that excessive through-thickness splitting (where the split spacing becomes small compared to the plastic-zone size) could be counterproductive to improving toughness.

5.3. Sources of weak interfaces

In commercial A1-Li alloys, the poor short- transverse properties in thick plate material result in part from the highly elongated and unrecrystallized grain structure (Fig. 3). Due to the large aspect ratios of the grains, a major portion of grain boundaries are aligned normal to the rolling plane; failures along these planes occur by brittle intergranular cracking, with less that 1% tensile elongation. In addition, deformation textures sustained during prior rolling and stretching [31-33], the presence of Cu- and Fe-rich intermetallics, lithium segregation to grain- boundary regions [48], heterogeneous precipitation [of equilibrium 6 (A1Li) particles or T l plates] along high-angle grain boundaries during artificial aging, and related solute-denuded 6'-PFZs, all contribute to the weakness of the grain-boundary interfaces. As a result, AI-Li alloys that display the best cryogenic toughness are generally the peak-aged and overaged tempers. Underaged microstructures, conversely, rarely show evidence of grain-boundary precipitation; the high-angle grain boundaries are thus com- paratively more resistant to short-transverse splitting,

50

40

30

20

"G

1o

• 2 9 8 K

- - ~ - - 7 7 K

0%% %..

i i i i

2 4 6 8 10

Thickness, B (ram/

Fig. 16. Variation in fracture toughness with sheet thickness (B) in near peak-aged 2090-T83 sheet at 298 and 77 K.


such that the toughness does not increase markedly at low temperature (Fig. 4).

6. CONCLUDING REMARKS

Apart from structural and mechanical-property anisotropy induced by the presence of aligned-weak interfaces, there are other limitations to delamination toughening. As noted above, in structural materials such as (monolithic or laminated) steels, aluminum and titanium alloys, laminating the microstructures beyond a critical dimension can lead to reduced toughness at ambient and elevated temperatures, despite enhanced low-temperature toughness compared to unlaminated structures. For example, attempts to decrease the ductile-brittle transition temperature (or improve low-temperature toughness) of mild steels by processing as laminates, are some- times accompanied by reductions in the upper-shelf toughness [20]. Thus, for crack-divider delamination toughening to be most effective, the laminate spacing must be comparable with the plastic-zone size, and less than or equal to the length of the split. In very strong solids, such as ceramics, layered granite rocks [49], and pyrolytic-carbon laminates [50], this is rarely achieved as, despite extensive splitting, the delami- nated segments remain in plane strain down to dimensions as small as a few microns. Although small contributions to toughness can arise from the energy expended in splitting, in these materials the orientations which invoke crack-arrester toughening clearly provide superior toughness. Finally, these benefits may not be realized during subcritical crack growth, even in ductile metals and alloys, because the lower stress levels are insufficient to cause laminar splitting along the weak interfaces [24].

7. CONCLUSIONS

Based on an experimental study of the micro- mechanisms influencing the cryogenic fracture- toughness behavior of commercial aluminum-lithium alloys, the following conclusions may be made:

1. All sheet and plate alloys tested show increased yield and tensile strength, tensile elongation and strain-hardening exponent with decrease in temperature from 298 to 77 K; the magnitude of the improvements is somewhat larger in thick-plate material.

2. Not all microstructures, specimen orientations and product forms, however, show improved fracture toughness at low temperatures, the effect being dependent on changes in the fracture mechanism and fracture mode (slant vs flat fracture) that occur over the temperature range.

3. For the long-transverse (L-T) orientation, microstructures that exhibit a transition in fracture mode or fracture mechanism, show reduced toughness at low temperatures. In 8090-T351, 8091-T351 and 2091-T8 thick-plate alloys, the reduction in

toughness is concurrent with a fracture-mechanism transition from (ductile) microvoid coalescence a t ' 298 K t.o (brittle) transgranular shear at 77 K; in 1.6 mm thin commercial 2090-T83 sheet alloy, it is concurrent with a fracture-mode transition from a 45 ° slant ("plane-stress") fracture at 298 K to a mixed mode (slant plus fiat) fracture at 77 K, even though the fracture mechanism (microvoid coalescence) remains unchanged.

4. Conversely, microstructures which exhibit no apparent change in microscopic fracture mechanism or macroscopic fracture mode display marked improvements in L-T toughness with decreasing temperature. In 2090-T81, 8090-T8, 8091-T8 and 2091-T351 thick-plate alloys that fail by transgranular shear at temperatures between 298 and 77 K, the increase in toughness is associated with an increasing tendency for short-transverse delamination cracking at low temperatures, which relaxes through-thickness constraint at the crack tip (crack-divider delamination toughening). Such through-thickness splitting divides the nominally plane-strain fracture process into a series of locally plane-stress (45%) slant failures. Additionally, toughness is found to increase with decreasing temperature for very thin 2090-T83 sheet, where a plane-stress fracture mode and a ductile microvoid coalescence fracture mechanism are preserved at all temperatures.

5. Consistent with the enhanced intergranular splitting at cryogenic temperatures, the short-transverse (S-L, S-T) toughness of all plate alloys tested shows a small decrease at low temperatures.

Acknowledgements--This work was supported by the Direc- tor, Office of Energy Research, Office of Basic Energy Sciences, Materials Sciences Division of the U.S. Depart- ment of Energy through Contract No. DE-AC03- 76SF00098. Thanks are due to Drs R. J. Bucci of Alcoa and W. E. Quist of Boeing for their advice and help in procuring materials, to Drs J. W. Morris, J. Glazer and R. H. Dauskardt for many stimulating (and often heated) discus- sions, to J. C. McNulty and J. E. Miles for experimental assistance, and to Madeleine Penton for her help in prepar- ing the manuscript.

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A P P E N D I X

Fracture Models

Stress -controlled fracture Micro-mechanical models for brittle transgranular cleavage and intergranular fracture in plane strain have been derived from a simple critical stress-controlled failure criterion [41, 43]. Specifically, the onset of crack extension at K = K~c is modelled in terms of a local criterion where the local maximum principal stress (fl) in the vicinity of the crack-tip exceeds a local fracture stress (f~') over a microscopically significant characteristic distance (r = l*), as depicted in Fig. Al(a). Assuming that the crack-tip stresses (fij), as a function of distance r from the tip, are described by the HRR singular solution [51, 52] for a nonlinear-elastic solid, such that as r ~ 0

F J 7 "/(l+'° o - ~ t r y / - - / ~/j(o, n) (A1)

Lf, E,/.rJ where J is the J-integral [53], given by K2/E ', E' is the elastic modulus appropriate to plane strain ( = E / ( 1 - r E ) , v is Poisson's ratio, try and Ey are the yield stress and yield strain, respectively, n is the strain-hardening coefficient, ~u(0, n) is a normalized stress function of 0 and n [54], and I n is a integration constant, weakly dependent upon n.

The onset of cracking, at K=KIc=(JIcE') 1/2, is thus consistent with fi exceeding f* over distance r = 1", such that

* (l+n)/2n (1 n)/ * 1/2 Klc~[{ff } /{fly} " 2'1{10 } . (A2)

Strain-controlled fracture (plane strain) Similarly, for ductile fracture under plane-strain con-

ditions, a (stress-modified) strain-controlled model [41,42] can be derived (Fig. Alb). Here, crack initiation, at J = Jic = K~c/E', is modelled in terms of the near-tip plastic strain (gp) exceeding a critical fracture strain (E*) over a characteristic dimension (r = r*). Assuming that the crack- tip strain distribution (E~) is given by the HRR solution [51, 52], i.e. as r - ,0

r J -]l/(l +.) %~Eyl ! ~,) (0, n) (A3)

kfyEylnr.J

2326 V E N K A T E S W A R A RAO and RITCHIE: CRYOGENIC F R A C T U R E - T O U G H N E S S OF A1-Li

(a) J =./~o '

5 - ~ ' ~ " ~ 3 - 5 ~ o

/ x,, .

I o II O.~) I i

0.02 0.03 I I

t ' - - % --'t

r / / (~ /o . ° )2

(b)

0 1 i ~ I

I.-%~ Fig. A1. Schematic idealization of microscopic plane-strain fracture criteria pertaining to (a) critical stress-controlled model for brittle fracture, and (b) stress-modified critical strain-controlled model for ductile fracture; after Refs

[41-43].

where g~ (0, n) is a normalized strain function of 0 add n [54], the fracture toughness for strain-controlled fracture is given " by

Kit ~ fl (E'aye~" r* )1/2 (A4)

where/7 = (l,/gu) ll2. The value of the local fracture strain E~ depends critically

on the degree of triaxiality or stress am/#, defined as the ratio of hydrostatic to equivalent stress. Circumferentially- notched tensile tests on AI Li (2090-T81) alloy plate [6] show the fracture strain appropriate to plane-strain conditions (O'm/# -+2), which is typical of the constraint in the vicinity of the crack tip, to be an order of magnitude smaller than the fracture strain for unconstrained (plane-stress) conditions (am/# --+ I/3). Such experimental results are consistent with the Rice and Tracey model [55] for ductile fracture which predicts the fracture strain to be inversely proportional to the exponental of 1.5 am/#, viz

E~ ~ ln(dolOp)l[0.28 exp(l.5 am/#)] (A5)

where dp/Dp is the ratio of spacing to diameter of the void initiating particles.

Plane-stress slant f racture For 100% slant fracture in thin sheet, where the stress-

state is essentially fully plane stress and the toughness decreases with decreasing thickness, the crack-tip displacement associated with crack extension has been modelled by Cottrell [46] in terms of the movement of screw dislocations on 45 ° planes ahead of the tip (Fig. A2). Following the treatment of Knot t [48], the critical strain energy release rate for such cumulative fracture is given by

2x/2 try B Gc-~ ~+~v) (A6)

which on simplifying and rewriting in terms of the stress intensity, yields the following expression for the fracture toughness

K c ~ (2EtryB) 112. (A7)

plane stress

Q

Shear produced by one dislocation MN

Movement of i - . screw dislocations

I'-,.

(. "" I i I

i i ~.• f y

i " ' . ~ . T I X

[

Fig. A2. Schematic idealization of the Cottrell model [461 for plane-stress 45 ° shear fracture; after Ref. [47].

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