Title of Publication Edited byTMS (The Minerals, Metal & Materials Society), Year

ENVIRONMENTAL INDUCED CRACKING in Al-Li-Cu-Mg-Zr ALLOYS of PEAK AGED and RETROGRESSED and REAGED TEMPERS

UNDER APPLIED POTENTIALS

K. S. Ghosh 1, K. Das 2 and U. K. Chatterjee 2

1 Department of Metallurgical and Materials Engineering, National Institute of Technology,Warangal – 506 004, India, ksghosh2001@yahoo.co.uk2 Department of Metallurgical and Materials Engineering, Indian Institute of Technology,Kharagpur – 721 302, India.

Keywords: Al-Li-Cu-Mg-Zr alloys, retrogression and reaging (RRA), anodic and cathodicpotentials, environmental induced cracking (EIC).

AbstractStress corrosion cracking (SCC) behaviour of 8090 and 1441 Al-Li-Cu-Mg-Zr alloys of peakaged T8, over aged T7 and retrogression and reaged (RRA) T77 tempers has been studied byslow strain rate test (SSRT) in 3.5 % NaCl and in 3.5 % NaCl + 0.1M LiCl + 0.7 % H2O2environments. The results show that a small addition of LiCl and H2O2 to NaCl makes the mediasusceptible to SCC. T8 temper has been found to be the most SCC susceptible, T7 temper theleast, and RRA T77 tempers are intermediate to T7 and T8 tempers. SSRT under appliedpotentials in 3.5 % NaCl solution shows that the alloy tempers are immune to SCC at freecorrosion potential and at small cathodic overvoltages, but at higher cathodic potentials,hydrogen embrittlement (HE) sets in. The applied anodic potentials are found to be severelydamaging, and local anodic dissolution (LAD) is believed to be associated mechanism for theSCC damage. RRA tempers are found to be more prone to HE compared to T8 temper and thishas been explained on the basis of microstructural variation observed by TEM, XRD and DSC.

IntroductionThe attractive combination of properties such as 5-10 % lower density, 15-20 % increase

of elastic modulus and 10 –15 % increase in specific strength, superior resistance to fatigue andcreep crack growth and higher cryogenic toughness has rendered the Al-Li alloys as candidatematerials for aerospace applications over the most widely used 2xxx and 7xxx series aluminiumalloys [1,2]. However, they have some drawbacks like the properties being strongly sensitive toprocessing conditions, unattractive fracture behaviour, high anisotropy of unrecrystallisedproducts as well as susceptibility to environmental induced cracking (EIC) [3].

The weight saving and increased stiffness benefits of the Al-Li alloys are useful to theaerospace vehicle designers provided they exhibit adequate resistance to EIC. Generally, the highstrength aluminum alloys of peak aged T8 tempers are most susceptible to SCC. Over aged T7temper has an acceptable SCC resistance [4,5,6], but is of lower strength requiring oversizecomponents. A novel heat treatment, retrogression and reaging (RRA) applied to the T8 temperhas found to have beneficial effect, as RRA treatment results in a microstructure approachingthat of theT7 temper while maintaining the T8 temper strength [7,8,9].

SCC behaviour of different aluminium alloys of various tempers has been studied byseveral authors [3-13]. Meletis [4] found in Al-2.9Cu-2.2Li-0.12Zr alloy that the degree of SCCresistance depends on the aging condition with over aging inducing susceptibility. In 8090 alloy,Gray [6] found the time to initiation of SCC in constant load testing has increased from the underaged (UA) to the peak aged (PA) and from the latter to the over aged (OA) condition. Lumsden

11

Degradation of Light Weight Alloys Edited by: David Shifler, Julie A. Christodoulou, James P. Moran, Airan Perez, Wenyue Zheng

TMS (The Minerals, Metals & Materials Society), 2007

and Allen [11] showed in their studies on 8090 alloy of UA, PA and OA conditions that theaging treatment did not change the SCC resistance. Alloy based on Al-Li-Cu-Zr and Al-Cu-Li-Zrsystems are known to exhibit maximum resistance to stress corrosion crack initiation in the nearT8 condition [10], whereas Al-Li-Cu-Mg-Zr alloys exhibited maximum resistance to crackinitiation in the T7 condition [6].

Al-Li alloys are also being considered for cryogenic tankage applications in thehypersonic and transatmospheric vehicles for their strategic importance. These space vehiclesare exposed to use liquid hydrogen as fuel, so the hydrogen embrittlement phenomenon in thesealloys for such applications is important to consider. The effect of cathodic charging of hydrogenon the mechanical properties of Al-Li alloys has also been studied by a few investigators and theobserved embrittlement was explained due to the formation of brittle hydrides [13]. However,studies on the effect of applied potentials on EIC of RRA T77 tempers of Al-Li-Cu-Mg-Zr seriesalloy are scanty. The present paper deals the study of SCC on 8090 and 1441 (Russian grade) Al-Li-Cu-Mg-Zr alloys of T8, and retrogression and reaging (RRA) T77 tempers in 3.5 % NaCl +0.1M LiCl + 0.7 % H2O2 solution and in 3.5 % NaCl solution under applied anodic and cathodicpotentials.

Experimental Details Materials

The 8090 and 1441 Al-Li-Cu-Mg-Zr alloys of T8 tempers were obtained in sheet formfrom Defence Metallurgical Research Laboratory (DMRL), Hyderabad, India and their chemicalcompositions (wt. %) are given in Table I. The as-received sheets were aged at 170 0C for 96hours to obtain T7 tempers. Figs. 1a and 1b, triplanar micrographs of the 8090 and 1441 alloys,exhibit the grain structures in the three orthogonal directions of the sheets. The grains of the8090 alloy are more or less equiaxed in the longitudinal (L) direction but are highly elongated inthe long transverse (LT) and in the short transverse (ST) directions. But, the grains of the 1441alloy are more or less equiaxed in all the three directions.

(a) (b)Fig. 1: Triplanar optical micrographs of (a) 8090 and (b) 1441 alloys.

Retrogression and Reaging (RRA) TreatmentThe T8 temper alloy sheets were subjected to retrogression treatment at temperatures near

and above the δ� (Al3Li) solvus line of Al-Li system for short duration, followed by reaging theretrogressed state to peak aged temper, which yielded four retrogression and reaged (RRA) T77tempers in each alloy. The details of RRA schedule are mentioned by the authors elsewhere[14,15]. However, in the present context, it is required to mention that for the SCC studies, RRAschedule was selected in such a way that reaging the retrogressed state has retained the initial T8temper strength and in the process, the total aging time of the RRA tempers become more thantwice to that of the conventional T8 temper, which is expected to evolve a microstructure of theRRA tempers approaching that of the T7 temper. The authors have characterized the

12

microstructural features of the alloys of T8, RRA T77 and T7 tempers by TEM, XRD, DSC,electrochemical polarization etc. techniques and discussed in elsewhere [16,17].

Table I: Chemical compositions (in wt %) of the 8090 and 1441 Al-Li-Cu-Mg-Zr alloys.Alloy Li Cu Mg Zr Fe Si Al8090 2.29 1.24 0.82 0.12 0.09 0.044 Balance1441 1.9 2.0 0.90 0.09 0.11 0.05 Balance

Slow Strain Rate Test (SSRT)For SSRT, tensile specimens of dimensions 25 mm gauge length, 4 mm width and 2.5

mm thickness for the 8090 alloy and 20 mm gauge length, 4 mm wide and 1.8 mm thickness forthe 1441 alloy were used, respectively. The surfaces of the gauge portions of the tensilespecimens were ground to 100 microns minimum depth so as to remove the lithium andmagnesium depleted and subsurface porosity zones, developed during solutionising carried out inair [18]. The gauge portions of the specimens were polished in emery papers upto 600 grits.

A CORTEST SSRT unit was used for testing at an initial strain rate of 6.0 X 10-6 s-1 forthe 8090 alloy and 5.5 X 10-6 s-1 for the 1441 alloy. Tests were carried out in air, 3.5 % NaCl, 3.5% NaCl + 0.1M LiCl + 0.7 % H2O2 and 3.5 % NaCl solution under applied anodic and cathodicpotentials. Tests were repeated to confirm the results. The plastic strain to fracture (εp), theductility ratio (DR) i.e. the ratio of plastic strain to fracture in environment to that in air (εenvn /εair), and the ratio of fracture energy (FR) i.e. the area under the plastic region of the stress straincurve in environment to that in air (Fenvn / Fair) are considered as the SCC susceptibility index.

For tests with applied potentials, an electrochemical cell was constructed by placing aplatinum wire mesh (acted as counter electrode) in a chamber containing test solution withinwhich the tensile specimen (acted as working electrode) was inserted. A salt bridge with aLuggin probe at one end was connected between the cell and a beaker containing saturated KClsolution and a saturated calomel electrode (SCE). Constant anodic and cathodic potentials wereapplied using an EG & G Princeton scanning potentiostat of model 362.

Results and DiscussionRRA and hardness

Figs. 2a and 2b show the variation of hardness with retrogression temperatures andreaging the retrogressed states to T8 aged tempers of the 8090 and 1441 alloys, respectively. Thefigures exhibit the characteristic RRA behaviour - a sharp decrease in hardness, attainment of aminimum, followed by a slight increase and then relatively constant with further retrogressiontime [14,19]; and reaging the retrogressed state (of retrogression time of minimum hardness) topeak aged duration resulted in attainment of initial hardness [15,20].

0 5 1 0 1 5 2 0 2 5 3 0 3 5 4 0 4 51 0 0

1 1 0

1 2 0

1 3 0

1 4 0

1 5 0

1 6 0

1 7 0

1 8 0

( a ) T i m e , m i n

Ha

rdn

es

s,

HV

10

R e t r o g r e s s io n a t 2 8 0 0 C R e t r o g r e s s io n a t 2 5 0 0 C R e t r o g r e s s io n a t 2 8 0 0 C a n d p e a k a g in g R e t r o g r e s s io n a t 2 5 0 0 C a n d p e a k a g in g

0 5 1 0 1 5 2 0 2 5 3 0 3 5 4 0 4 5 5 0

1 0 0

1 1 0

1 2 0

1 3 0

1 4 0

1 5 0

1 6 0

( b )

Ha

rdn

es

s,

HV

10

R e t r o g r e s s io n a t 2 7 0 0 C R e t r o g r e s s io n a t 2 7 0 0 C a n d p e a k a g in g R e t r o g r e s s io n a t 2 3 0 0 C R e t r o g r e s s io n a t 2 3 0 0 C a n d p e a k a g in g

T i m e , m i n

Fig. 2:Variation of hardness with retrogression time at different retrogression temperature andreaging to peak aged tempers of (a) 8090 and (b)1441 alloys.

13

Microstructural FeaturesTEM micrographs: Fig. 3, a few representative TEM micrographs of the 8090 alloy, shows anumerous matrix strengthening δ� (Al3Li) precipitates in T8 temper (Fig. 3a), equilibrium δ (AlLi) phase along the grain and subgrain boundaries in the RRA 8090R280IA temper, and auniform distribution of S� (Al2CuMg) and T1 (Al2CuLi) precipitates in the RRA 8090R280DAtemper (Fig. 3c). Authors have discussed the microstructural features of the T8, RRA and T7tempers of the alloys, and concluded that RRA tempers have more or less same amounts and sizeof δ� precipitates, lower dislocation density, and higher amounts of equilibrium δ and T1 and S�

precipitates compared with that in the T8 temper[14-16] and similar observations by others [8,9].

(a) (b) (c)Fig. 3: TEM microstructures of (a) 8090-T8 shows numerous δ� precipitate, (b) RRA8090R280IA temper display equilibrium δ phase along the grain and subgrain boundaries, and(c) RRA 8090R280DA with uniform distribution of S� and T1 phases.

DSC studies: Figs. 4a and 4b show the DSC thermograms of the 8090 and 1441 alloys in theirT8, RRA and T7 tempers, respectively. The exothermic and endothermic peaks representingprecipitations sequence are well established [17,21] and the peak D region represents theprecipitation of S�� T1 and δ phases. Comparing the area under the peak region D (which is anindication of proportion of precipitation [22]) and increase intensity of the D2 peak of the T7temper, it can be inferred that the RRA and T7 tempers do contain higher amounts of S�� T1 and δ phases with respect to that of the T8 tempers.

5 0 1 0 0 1 5 0 2 0 0 2 5 0 3 0 0 3 5 0 4 0 0 4 5 0 5 0 0- 0 .9- 0 .8- 0 .7- 0 .6- 0 .5- 0 .4- 0 .3- 0 .2- 0 .10 .00 .10 .20 .3

( a )

BA E

D 1

D2

D

C

8 0 9 0 - T 8 8 0 9 0 - T 7 8 0 9 0 R 2 8 0 D A

He

t fl

ow

(m

ca

l/s

ec

)

T e m p e r a tu r e , 0 C5 0 1 0 0 1 5 0 2 0 0 2 5 0 3 0 0 3 5 0 4 0 0 4 5 0 5 0 0

- 0 .8

- 0 .7

- 0 .6

- 0 .5

- 0 .4

- 0 .3

- 0 .2

- 0 .1

0 .0

( b )

EC

D 2D 1

He

ta f

low

(m

ca

l/s

ec

)

T e m p e r a t u r e , 0 C

1 4 4 1 -T 8 1 4 4 1 -T 7 1 4 4 1 R 2 3 0 D A

Fig. 4: DSC thermograms of alloys (a) 8090 of T8, RRA 8090R280DA and T7 and (b) 1441 ofT8, RRA 1441R230DA and T7 tempers, at a heating rate of 10 0 C/min.

SCC behaviour of T8 and RRA tempersFigs. 5(a-b) and 5(c-d) represent stress strain curves of T8 and RRA tempers of the

alloys, tested at slow strain rate in air, in 3.5 % NaCl + 0.1M LiCl + 0.7 % H2O2 solution, and 14

under applied anodic and cathodic potentials. Tables II and III give the average values ductilityratio (DR) and fracture energy values of alloys of various tempers, in different media, at ananodic and a cathodic potential.

The stress strain curves (Fig. 5) and Tables II and III indicate that the reduction of plasticstrain to fracture is greatest when an applied potential is applied in any solutions with saltsupporting electrolyte. Fig. 5 also exhibits that the addition of a small amount of LiCl and H2O2to 3.5 % NaCl solution makes the media conducive to SCC, and the applied cathodic potentialhas also affected to SCC damage. Tables II and III and the stress strain curves (Fig. 5) alsoindicate that the RRA tempers exhibit higher SCC resistance than to that in the T8 tempers of thealloys in 3.5 % + 0.1M LiCl + 0.7 % H2O2 solution. Hence, the RRA treatment resulted in animprovement of SCC resistance and similar results are also reported by others [5,8,9,23,24].

The applied stress during SSRT assists in rupturing the passive film and leads to the localdissolution of the exposed area as it acts as a small anode in contact with a large cathode i.e. thefilm surface. The film rupture exposes an area containing precipitates and precipitate free zonesthat are electrochemically different providing local galvanic cells.

The microstructure of the T8 temper comprises the precipitations of δ�, S�, T1, β� phaseswithin the matrix, and δ, T1 and β� phases at the grain boundaries (Figs. 3 and 4) [16,17]. Theinitiation of SC cracks is believed to be caused by the preferential dissolution of the anodic grainboundary and sub-grain boundary δ and T1 precipitate by forming a lots of galvanic cells. Thenucleation of crack has also occurred from the pit base. Further, the crack initiation is assisted bythe strain localization arising due to coarse planar slip generated by the coherent δ' precipitatesand equlibrium δ at the grain boundary. Thus, an anodic dissolution promoted intergranular SCC

0 1 2 3 4 5 6 7 85 0

1 0 0

1 5 0

2 0 0

2 5 0

3 0 0

3 5 0

4 0 0

4 5 0

( a ) S t r a in , %

Str

es

s,

MP

a

A i r 3 . 5 % N a C l 3 . 5 % N a C l + 0 . 1 M L i C l + 0 . 7 % H 2 O 2 A n o d i c p o t e n t i a l - 6 7 5 m V ( S C E ) C a t h o d i c p o t e n t i a l - 1 5 5 0 m V ( S C E )

0 1 2 3 4 5 6 7 8 9 1 0 1 15 0

1 0 0

1 5 0

2 0 0

2 5 0

3 0 0

3 5 0

4 0 0

4 5 0

( b )S t r a in , %

Str

es

s,

MP

a

A i r 3 .5 % N a C l 3 .5 % N a C l + 0 .1 M L iC l + 0 .7 % H 2 O 2 A n o d ic p o t e n t ia l - 5 2 5 m V ( S C E ) C a t h o d ic p o t e n t ia l - 1 5 5 0 m V ( S C E )

0 1 2 3 4 5 6 7 8 9 1 05 0

1 0 0

1 5 0

2 0 0

2 5 0

3 0 0

3 5 0

4 0 0

4 5 0

(c ) S tr a in , %

Str

es

s,

MP

a

A ir 3 .5 % N a C l 3 .5 % N a C l + 0 .1 M L iC l + 0 .7 % H 2O 2 A n o d ic p o te n t ia l -5 7 5 m V (S C E ) C a th o d ic p o te n t ia l -1 5 5 0 m V (S C E )

0 1 2 3 4 5 6 7 8 9 1 0 1 1 1 25 0

1 0 0

1 5 0

2 0 0

2 5 0

3 0 0

3 5 0

4 0 0

4 5 0

(d )S tra in , %

Str

es

s,

MP

a

A ir 3 .5 % N a C l 3 .5 % N a C l + 0 .1 M L iC l + 0 .7 % H 2O 2 A n o d ic p o te n t ia l -5 6 5 m V (S C E ) C a th o d ic p o te n t ia l -1 5 5 0 m V (S C E )

Fig. 5: Representative stress strain curves of (a) 8090-T8, (b) 8090R280IA, (c) 1441-T8 and (d)1441R230DA tempers in different SCC media, and in applied cathodic and anodic potentials.

15

mechanism seems to be operative in the alloy and intergranular cracking is evident from thestress corrosion (SC) crack path and SEM fractographs, discussed in subsequent sections.

The SCC process in the RRA tempers is believed to be the same to that in the T8 temper.But, the improved SCC resistance of RRA tempers can be explained based on the microstructuralchanges that have resulted in the formation of additional equilibriumδ and T1 phases (Figs. 3 and4) [14-17]. The more dispersion of S' and T1 phase within the matrix, and equilibrium δ phasealong the grain boundaries and sub-grain boundaries, reduce the potential difference for theformation of microscopic galvanic cells. Therefore, the tendency and propensity of intergranulardissolution is largely reduced resulting in an improvement of resistance to SC crack initiationand propagation.

Table II: Mechanical properties of the 8090 alloy of different tempers in different media.Alloy

Temper3.5 % NaCl 3.5 % NaCl +

0.1MLiCl + 0.7% H202

Cathodicpotential,

-1550 mV (SCE)

Anodicpotential,

-525 mV (SCE)DR

(εenv/(εair)

Fractureenergy,J/cm3

DR(εenv/(εair)

Fractureenergy,J/cm3

DR(εenv/(εair)

Fractureenergy,J/cm3

DR(εenv/(εair)

Fractureenergy,J/cm3

8090-T8 0.78 13.40 0.62 12.50 0.83 12.44 0.32 5.17

8090R280IA 0.89 23.97 0.79 17.10 0.71 13.50 0.22 4.04

8090R280DA 0.91 23.89 0.85 17.50 0.70 16.30 0.20 5.70

8090R250IA 0.86 22.94 0.82 14.11 0.64 12.50 0.17 4.56

8090R250DA 0.94 28.56 0.74 16.80 0.79 17.00 0.17 3.50

Table III: Mechanical properties of the 1441 alloy of different tempers in different media.Alloy

Temper3.5 % NaCl 3.5 % NaCl +

0.1MLiCl + 0.7% H202

Cathodicpotential,-1550 mV

(SCE)

Anodic potential,-575 mV (SCE)

DR(εenv/(εair)

Fractureenergy,J/cm3

DR(εenv/(εair)

Fractureenergy,J/cm3

DR(εenv/(εair)

Fractureenergy,J/cm3

DR(εenv/(εair)

Fractureenergy,J/cm3

1441-T8 0.93 24.05 0.36 7.26 0.96 24.2 0.29 5.12

1441R270IA 0.88 24.40 0.54 10.50 0.88 17.64 0.35 4.04

1441R270DA 0.87 23.63 0.71 15.00 0.75 18.08 0.40 5.70

1441R230IA 0.94 24.60 0.59 12.00 0.83 21.76 0.35 4.56

1441R230DA 1.00 27.85 0.74 16.00 0.87 21073 0.17 3.20

Effect of applied potentialsFigs. 6(a-e) and 7(a-e) show the average values of ductility ratio (DR) vs. applied

potentials of T8 and RRA tempers of the 8090 and 1441 alloys, tested at slow strain rate in 3.5 %NaCl solution, respectively. Figs. 6 and 7 exhibit that the applied anodic potentials have a drasticdecreasing effect on ductility values in all the tempers. Under cathodic potentials, within thecathodic overvoltage of the order of approximately 200 mV (SCE), the DR of all the tempersremained unaffected. However, with further decrease of cathodic potential upto –1500 mV

16

(SCE), there are decreases of ductility ratios in the RRA tempers of both the alloys. The DRvalues at -1550 mV (SCE) of all the RRA tempers are lower compared to the T8 tempers.

The applied potential versus ductility ratio curves (Figs. 6 and 7) can be divided in threeregions i.e. under anodic overvoltage (region I), under low (~ 200 to 400 mV) cathodicovervoltage (region II) and high (~ 400 to 800 mV) cathidic overvoltage (region III). In theregion I of anodic overvoltage, metal ion dissolution is the most prominent reaction leading toSCC with a drastic drop in ductility ratio values. In the cathodic overvoltage regions II and III,hydrogen evolution reaction at the specimen surface predominates. So, there will be hydrogeningress into the metal matrix through the defects and the oxide layers. The amount of hydrogendiffusion for a given alloy structure and oxide layer depends on the relative hydrogen content atthe surface. In the region (III) under higher cathodic overvoltage, vigorous hydrogen evolutionreaction occurs. But, in the region (II) with lower cathodic overvoltage, the hydrogen evolutionreaction obviously is rather slow. The high value of ductility ratio around the OCP is indicativeof the absence of stress corrosion cracking (SCC) or hydrogen embrittlement (HE). The amountof hydrogen diffusion into the metal matrix will be more with higher cathodic potentialcompared to the hydrogen diffusion with lower cathodic potential. The low ductility ratio inregion (III) with higher cathodic potential reflects that hydrogen plays an important role inreducing ductility or, in other words, hydrogen embrittlement sets in [25]

The hydrogen embrittlement in these alloys could be due to several features. Thegeneration of hydrogen on the specimen surfaces during applied cathodic potential and itsdiffusion into the material results in the formation of a brittle, coherent hydride phase (LiAlH4),which initiates microcracks. The hydride phase is associated with the δ phase (AlLi) accordingto the reaction (1) that exists preferentially in grain and subgrain boundaries.

AlLi (δ) + 4H � LiAlH4 (1)Thus, SCC at this high cathodic potential is assisted by hydrogen embrittlement reaction due tohydride formation and its brittleness [13]. The amount of δ phase present is more in the RRAtempers compared to the T8 tempers. So, the tendency of formation of brittle hydride phaseobviously will be more in the RRA tempers, as reflected with the DR drop (Figs. 6 and 7).

According to Meletis [26] hydrogen ions are discharged on T1 precipitates or on copperdeposits in the crack tip area produced by the dissolution of grain boundary T1 phase, and atomichydrogen diffuses fast along grain boundaries, especially the ones which are vertical to the stressaxis. Thus, crack initiation occurs at these favourably positioned grain boundaries, due to HE andor dissolution of T1 and then the crack propagates when the hydrogen concentration in front ofthe crack tip reaches a critical level. Since RRA treatments have also resulted in a higher amountof T1 precipitates [16,17] at the grain boundaries as well as within the grain, the propensity of HEwill be more in the RRA tempers than that in the T8 tempers, where the critical hydrogenconcentration is achieved, at higher applied cathodic potentials, also reported in literature [25].

Stress corrosion crackFig. 8(a-c) exhibits the stress corrosion crack initiation and propagation path in the 1441

alloy. The SC crack initiates in the denuded region at the base of a pit (Fig. 8a), the propagationis intergranular (Fig. 8b) in 3.5 % NaCl + 0.1M LiCl + 0. 7 % H2O2, and also intergranularpropagation (Fig. 8c) under applied anodic potential.

SEM FractographsFig. 9(a-i) shows a few representative SEM fractographs of the alloys of various tempers

tested at slow strain rates in air, SCC media and under applied potentials. Fig. 9(a-d), thefractographs of the 8090 alloy shows as a whole a fibrous appearance - characteristic of thematerials that possess a highly textured structure, a type of fracture known as flutting. Fig. 9(e-i),the fractographs of the 1441 alloy reveals mostly intergranular and sub-intergranular cracking inpresence of a few dimples, irrespective of the tempers and media of testing. The intergranular

17

nature of cracking in the alloys is associated with strain localization arising from the shearing ofcoherent ordered δ' precipitates by moving dislocations and the subsequent strain concentrationat the grain boundaries. The dissolution of the precipitates at the grain boundaries and sub-grainboundaries in the SCC media and under applied potential has further contributed to the initiationof intergranular cracks. So, the intergranularity appears to be somewhat more in the specimenstested under applied anodic potential (Figs. 9c, 9d, 9h and 9i).

-1500 -1200 -900 -600

0.2

0.4

0.6

0.8

1.0

-1500 -1200 -900 -600

0.2

0.4

0.6

0.8

1.0

OCV = -725 mV

Duc

tility

rat

io

Duc

tility

rat

io

Duc

tility

rat

ioD

uctil

ity r

atio

Duc

tility

rat

io

-1500 -1200 -900 -600

0.2

0.4

0.6

0.8

1.0

OCV = -740 mV

(e)

(d)(c)

(b)(a)

Fig. 6: Ductility ratio of 8090 alloy of

(a) T8, (b) 8090R280IA, (c) 8090R280DA,

(d) 8090R250IA and (e) 8090R250DA

tempers with applied potentials in

3.5 % NaCl solution.

-1500 -1200 -900 -600

0.2

0.4

0.6

0.8

1.0

OCV = -731 mV

-1500 -1200 -900 -600

0.2

0.4

0.6

0.8

1.0

OCV = -734 mV

Applied potentials, mV (SCE)

Applied potentials, mV (SCE)Applied potentials, mV (SCE)

Applied potentials, mV (SCE)

Applied potentials, mV (SCE)

OCV = -736 mV

-1500 -1200 -900 -600

0.2

0.4

0.6

0.8

1.0

-1500 -1200 -900 -600

0.2

0.4

0.6

0.8

1.0

(a)OCV = -732 mV

Duct

ility

ratio

Duct

ility

ratio

Duct

ility

ratio

Duct

ility

ratio

(b)

Fig. 7: Ductility ratio of 1441 alloy of

(a) T8, (b)1441R270IA, (C) 1441R270DA,

(d) 1441R230IA and (e) 1441R230DA

tempers with applied potentials in

3.5% NaCl solution.

-1500 -1200 -900 -600

0.4

0.6

0.8

1.0

1.2

OCV = -718 mV

(c) -1500 -1200 -900 -600

0.2

0.4

0.6

0.8

1.0

OCV = -727 mV

Applied potentials, mV (SCE)

Applied potentials, mV (SCE) Applied potentials, mV (SCE)

(d)Applied potentials, mV (SCE)

-1500 -1200 -900 -600

0.2

0.4

0.6

0.8

1.0

OCV = -728 mV

OCV = -697 mVDuct

ility

ratio

(e)Applied potentials, mV (SCE)

Fractographs (Fig. 9b and 9f) of the RRA tempers have exhibited more dimples withinthe intergranularly separated grains compared to that observed in the T8 tempers (Fig. 9a and9e). This is because that the RRA treatments resulted in lowering of dislocation densities andalso enhanced precipitation of semi-coherent S' and T1 in the matrix which tends to homogenizeslip localization by promoting cross-slip. This is also reflected with higher ductility values, testedin air of the RRA tempers (Figs. 5b and 5d) compared with that of (Figs. 5a and 5c).

50 μm

(a)

50 μm

(b) (c)Fig. 8: (a) SC crack initiation from the base of a pit and from denuded region, (b-c) propagationis intergranular in 3.5 % NaCl + 0.1MLiCl + 0. 7% H2O2 and in applied anodic potential.

Conclusions1. Retrogression treatment subjected to T8 tempers of the 8090 and 1441 alloys and immediatelyreaging the retrogressed states to T8 duration caused in retaining the initial T8 hardness.2. Slow strain rate tests show that the addition of small amounts of LiCl and H2O2 to NaCl,makes the media conducive to SCC by maintaining a condition of passivasion required for SCC.

50 μm

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3. RRA treatment has resulted in an improvement of SCC resistance because of decreasedgalvanic effect caused by the enhanced precipitations of δ and T1 phases. But, the SCCmechanism is believed to be same in the RRA and T8 tempers.4. Applied anodic potential has enhanced SCC in both the alloys, and at higher applied cathodicpotentials hydrogen embrittlement has set in. RRA tempers have been found to be more prone tohydrogen embrittlement compared to the T8 temper, which might be due to the presence ofhigher amounts of δ phase, enhancing the formation of brittle hydride phase, an or the presenceof higher amounts of T1 phase providing sites for cathodic depolarization.5. Stress corrosion crack initiation is from the base of the pit and denuded regions but the crackpropagation is predominantly intergranular. SEM fractigraphs reveling flutting type of fracture inthe 8090 alloy and integranular cracking in the 1441 alloy.

(a) (b) (c)

(d) (e) (f)

(g) (h) (i)Fig. 9: SEM fractographs of (a) 8090-T8 in air, (b) 8090R280IA in air, (c) 8090R250IA in ananodic potential, (d) 8090R280IA in an applied anodic potential, (e) 1441- T8 in air, (f)1441R230IA in air, (g) 1441R230DA in 3.5 % NaCl + 0.1M LiCl + 0.7 % H2O2 solution and (h-i) 1441-T8 in an applied anodic potential.

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