Effects of Hydrogen on Mechanical Propertiesand Fracture Mechanism of 8090 AI-Li Alloy

LIAN CHEN, WENXIU CHEN, ZHONGHAO LIU, YUXIA SHAO, and ZHUANGQI HU

The effects o f hydrogen and strain rate on the mechanical properties and fracture mechanismo f 8090 A1-Li alloy under electrochemical charging conditions have been studied. Experimentalresults demonstrate that the tensile strength [ultimate tensile strength (UTS) and yield strength(YS)] and plasticity [reduction o f area (RA) and elongation (EL)] drop linearly with the decreaseo f strain rate. The charged hydrogen increases the tensile strength but markedly impair theplasticity. The susceptibility o f hydrogen embrittlement increases with the decreases o f strainrate, and the susceptibility o f the charged specimens was larger than that o f the uncharged onesover the strain-rate range. Observation by scanning electron microscope (SEM) reveals that thecharged hydrogen enhances intergranular delamination cracking on the fracture surface. Thefracture model o f charged specimens at low strain rates (~ < 3.4 x 10-4/s) is grain boundarybrittle fracture (GBBF), while that o f other conditions is grain boundary ductile fracture (GBDF).Secondary ion mass spectroscopy (SIMS) study shows that the atomic binding energy o f A1 andLi in the alloy decreased after hydrogen charging, and the atomic binding energy drop o f theformer is more than the latter. In this article, the hydrogen transport through the mobile dis-location mechanism o f hydrogen-induced fracture and the hydrogen effect on atomic bindingenergy were also discussed in detail.

I. INTRODUCTION

ALUMINUM-LITHIUM alloy is a new kind of struc-tural material. Its density is less than that o f conventionalaluminum alloy, and its strength and modulus are higherthan those o f the latter. It has great promise fo r potentialapplications in the aerospace and weapon industries.l 1-7]Nevertheless, the effects of hydrogen on the mechanicalproperties and fracture behavior o f A1-Li alloy have be-come so important that they need to be studied urgently.Although in recent years a few articles have been pub-lished on hydrogen embrittlement of 8090 and 2090 A1-Lialloys, I8-~l~ many problems remain to be solved. Thus,in this article, the effects o f hydrogen and strain rate onthe mechanical properties and fracture mechanism o f 8090A1-Li alloy are studied.

II. EXPERIMENTAL P R O C E D U R E S

The chemical composition o f 8090 A1-Li al loy is listedin Table I. The alloy was vacuum melted in an inductionfurnace and cast in a high-purity argon atmosphere. Theingot was annealed at 783 K for 12 hours, and then thesurface layer was machined away . Af t e rbeing preheatedat 783 K for 2 hours, the ingot was hot rolled to platesof 2-mm thickness.

The plate was solution treated at 803 K for 1 hour andwater quenched at room temperature. All specimens wereaged at 463 K for 16 hours in o rde r to obtain a micro-structure corresponding to the peak-aged condition. Thegrain structure o f the plate was partially recrystallized

LIAN CHEN, Professor, WENXIU CHEN, Professor, ZHONGHAOLIU, Assistant Professor, YUXIA SHAO, Assistant Professor, andZHUANGQI HU, Professor, are with the Institute of Metal Research,Academia Sinica, Shenyang 110015, People's Republic of China.

Manuscript submitted April 23 , 1992.

with larger pancake-shaped grains. Owing to the pres-ence o f pancake-shaped grains, tensile specimens weretaken along the long transverse directions and perpen-dicular to the rolling direction. The tensile specimenshad a gauge length o f 26 mm and width o f 6 mm.

All o f the specimens were carefully polished to re-move the surface damage. After polishing, some o f thesespecimens were cathodically hydrogen charged in a 0.1N NaOH solution containing 250 mg A s 2 0 3 p e r liter.The charging current density was 100 A / m e, and the hy-drogen charging time was 6 hours. At room temperature(293 K), the tensile test was carried out on a ShimadzuAutograph Ag-6000A testing machine to fracture withthree kinds of strain rates (~): 3,4 x 10-3/s , 3.4 x 10 4/s, and 3.4 x 10 5/s. The fracture morphology was ob-served by a Shimadzu EPM-810 electron microscope.

Commercial secondary ion mass spectroscopy (SIMS)equipment o f "second generation," KYKY LT-1A, wasused to detect the change o f binding energy o f aluminumatom before and af te r hydrogen charging. A schematicdiagram o f LT-1A type o f SIMS is shown in Figure 1,in which an energy fil ter grid was inserted between thecollection slit and transfer electrode. Changing the gridpotential, beginning with the lowest at 5 V to the high-est, can measure the secondary ion beam intensity andretarding potential. The curve o f the secondary ion beamvs the retarding potential was recorded by x-y recorder.With the increase o f grid potential, the number o f sec-ondary ions passing through the grid potential barr ierde-crease. The middle part o f the curve is nearly linear, andit will shift from the left to the right if the binding energybetween aluminum atoms and lithium atoms decreases;that is, more 27A1+ and 7Li+ will be sputtered. Therefore,the retarding potential o f hydrogen-charged specimens ishigher than that o f uncharged ones, The parameters o fSIMS are as follows: the primary beam is Ar+; the ac-celerating voltage is 15.5 kV; the diameter and the cur-rent o f the beam on the sample surface are 50 to 6 0 / z m

METALLURGICAL TRANSACTIONS A VOLUME 24A, JUNE 1993--1355

T a b l e I . C h e m i c a l C o m p o s i t i o n o f 8 0 9 0 A I - L i A l l o y W e i g h t P e r c e n t

Li Cu Mg Zr Fe Si Na K Ca A1

2.59 1.30 1.10 0.13 0.012 <0.01 0.002 0.005 0.002 bal.

cylindrical condenserPrimary ion

%Repel l el"

/~¢¢'/ %_ % Ispeclmeal % * .

~C O ' . .

homogeneous

magnetic field

Fig. 1--Schematic diagram of LT-1A type of S IMS.

and 0.5 /xA, respectively; and the ultimate pressure inthe sample chamber is about 5 x 10-5 Pa.

III. EXPERIMENTAL R E S U L T S

A. Mechanical Properties and HydrogenEmbrittlement Index

The uniaxial tensile results are shown in Figures 2 and3. It can he clearly seen that the tensile strengths [ulti-mate tensile strength (UTS) and yield strength (YS)] o f

5 0 0

~ 4 0

t--IlJk.

3 0Itl1/1c-(DF-

2 0

10

0 -

0 -

O -

UTS YSI g O

J •

O . - - - - - - -t3-

unchargedc h a r g e d

j D- - O ~ , - l i d

- I 1 ~

Test tempera?ure Z98K

--@

I I- $ -& "3

o 1 0 10

Strain R a t e ( s e e " )

Fig. 2 - - E f f e c t of hydrogen and strain rate on the tensile strength of8090 A1-Li alloy.

hydrogen-charged specimens drop linearly with thedecrease o f strain rate, while those o f uncharged onesdecrease only a little. The tensile strengths o f hydrogen-charged specimens are higher than those o f unchargedones, and this is in accordance with the results of Binfieldet al.ll2] It is interesting to note that the plasticity [re-duction o f area (RA) and elongation (EL)] o f chargedspecimens is decreased greatly, while those o f un-charged specimens is also decreased with strain rate. TheRA o f the uncharged specimens decreases from 5.0 to4.0 pct over the strain-rate range. Clearly, the hydrogen-charged specimens were more dependent on strain rate.

The hydrogen-induced brittleness susceptibility o f thealloy can be expressed by the hydrogen embrittlementindex IH, as IH = [(~00 -- CH)/qJO] × 100 pct, where ¢oand ~0. are RA o f uncharged and charged specimens,respectively. The calculation results, based on the datashown in Figure 3 and the above equation, are shown inFigure 4. The u p p e r curve indicates the effect o f strainrate and hydrogen charging on the hydrogen embrittle-ment index, while the other curve represents the effecto f strain rate on the hydrogen embrittlement index. It canbe clearly seen that the hydrogen embrittlement indexincreases with the decrease o f strain ra te , and the em-brittlement index o f the charged specimens is larger thanthat o f uncharged ones over the strain-rate range. There-fore , this indicates that the al loy is susceptible to hydro-gen embrittlement that can be caused by cathodic hydrogencharging.

B . Fractography

Scanning electron microscope (SEM)observation showsthat there are many secondary cracks on the surface o fall o f the tensile specimens. These secondary cracks aredelamination crack o f along pancake-shaped grainboundaries in the short transverse direction. This resultis in accordance with the results o f other researchersJ 9']3]However, there are more obvious delamination crackson the fracture surfaces o f hydrogen-charged tensilespecimens than those o f uncharged ones (Figure 5). Thehigher magnification observation is shown in Figure 6.Many microvoids exis t on intergranular fractured sur-faces o f uncharged specimens, and the number and sizeo f microvoids seem to be lowered with the decrease o fstrain ra te . Deformed microvoids on intergranular frac-ture surfaces are tested at 3.4 × 10-3/s strain rate, whichis similar with that o f hydrogen uncharged ones, whereasthose at strain rates 3.4 x 10-4/s and 3.4 × 10-5/s aresmooth and flat intergranular fracture o f charged speci-mens. The former is a grain boundary ductile fracture(GBDF), while the latter is a grain boundary brittle f rac-ture (GBBF).

C. Atomic Binding Energy

The relationship between intensity o f secondary ion27A1+ and 7Li+ and retarding potential is shown in

1356--VOLUME 24A, JUNE 1993 METALLURGICAL TRANSACTIONS A

6 . 0

5 , 0

¢ .

~o 4.0,.Co

3 . 0

RA EL© e u n c h a r g e d

• • c h a r g e d ID

T e s t t e m p e r a t u r e 2 9 8 K

2.0 I II d5 16" I cfs

S t r a i n R a t e ( s e e " )

Fig. 3 - - E f f e c t of hydrogen and strain rate on the plasticity of 8090A1-Li alloy.

5O

z,O

30

2o

T e s t t e m p e r a t u r e 1 5 I K

01 I I ~-~_10" 10.4 10"

STRAIN RATE(sec")Fig. 4 - - 8090 AI-Li al loy hydrogen embrittlement index as a functionof hydrogen and strain rate.

Figures 7 and 8, respectively. It can be seen that for thesame retarding potential, the intensity o f secondary ion27A1+ and 7Li+ o f hydrogen-charged specimens increasesquite little as compared with that o f uncharged ones, butthe shape o f the curves remains the same. This impliesthat hydrogen reduces the atomic binding energies o f A1and Li. By means o f l inear regression method, the inter-secting points o f the two curves extrapolated on the re-tarding potential coordinate in Figures 7 and 8 weredetermined. The difference o f the two crossing pointswas found to be 3.2 and 2.3 V for aluminum and lithiumelements, respectively, which means that, after hydro-gen charging, the decreases in atomic binding energiesof aluminum and lithium are 3.2 and 2.3 eV, respectively.

IV. DISCUSSION

A . Hydrogen Transport by Dislocations

Bastien and AzouD4I first suggested that mobile dis-locations can transport hydrogen, and numerous inves-tigators have proven that plastic deformation can causethe movement o f dislocations carrying hydrogenJ 15-mThe result o f this work also indicates that mobile dis-locations can capture hydrogen atoms and carry them,moving away at slow strain rate. In terms o f transportinghydrogen by dislocation, now there are two models: oneis the Cottrell atmosphere model[]9] and the other is thedislocation core atmosphere model, t2°] Which one is morepractical? We consider the latter model to be better forthe experimental results, because the cores o f disloca-tions are heavily distorted regions in crystal lattice, whichcauses a strong strain field. When hydrogen atoms fillthis region, they will relax the strain field and result ina decrease o f distortion energy; therefore, dislocation coresare strong hydrogen traps and are 3 to 5 times strongerthan those o f the Cottrell atmosphere, [2°1 so that hydro-gen atoms along dislocation lines will be trapped by dis-location cores and then carried away to surface a n d / o rgrain boundaries and interfaces o f precipitates. Fuentes-Samaniego and Hirth12t] derived the critical velocity V~for breakaway o f core atmosphere as

Vc = 3 . 2 ( D H / a ) [ I ]

where D , is the diffusivity o f hydrogen atom in matrix(in the case o f A1-Li alloy, it is 2.6 × 10-14mZ/stZe]) anda is the radius o f a dislocation core, taken as b , the Burgersvector. The critical strain rate ~c o f specimens is givenas

i c = 0 , , b . V,. [2]

where Po is mobile dislocation density, taken as 106/c m 2

By combining Eqs. [1] and [2]

ec = 3.2 (pn'DH) = 8.3 X 10-4/s [3]

In this work, strain rates gc (3.4 x 10-4/S and 3.4 x10-5/s) are less than the critical strain rate ic (8.3 x10-4), i.e., the mobile velocity V for drawing the dis-locations less than the critical velocity Vc for breakawayo f the core atmosphere. Therefore, in this case, mobile

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Fig. 5--Fracture surfaces in low magnification (SEM): (a) uncharged, ~ = 3 .4 x 10 4/s, (b) charged, ~ = 3 .4 x 10-5/s; (c) uncharged, g =3.4 x 10 3/s; and (d) charged, g = 3 .4 x I0 4/s.

dislocations play a main role in hydrogen transportation;but as to the case o f higher strain rate (3.4 × 10 3/s) ,the hydrogen bulk diffusion may be the effective methodo f hydrogen transportation.

B. Mechanism o f Hydrogen-Induced Fracture

According to the test results, the fracture mode o f ten-sile test is delamination fracture. The fracture surfacesexhibit featureless and secondary cracks. The cracks cor-respond to the pancake-shaped grain boundary delami-nation cracks along the short transverse direction. As statedabove, the fracture surfaces o f uncharged specimens atthree strain rates and hydrogen-charged ones at a strainrate o f 3.4 × 10-3/s are all G B D F , mainly due to theprecipitates along the grain boundary. In the hydrogenuncharged case, because o f the low hydrogen content inspecimens and the fast strain rate, although mobile dis-locations carrying hydrogen can cross the precipitate freezone (PFZ) under applied stress, no heavy accumulationon grain boundary occurs, so cracks nucleate in the vi-cinity o f precipitates on grain boundaries to cause GBDF.This is consistent with the results o f other researchers.j23]Nevertheless, as in the case o f hydrogen charging, thehydrogen content increases dramatically. When the strainrate is slow enough, the dislocations interact with hy-drogen atoms to form a core atmosphere, and hydrogenatoms are carried into the PFZ to cause cracking. Thisenhances dislocation slip on the slip plane in the PFZand usually results in a microscopically smooth GB frac-ture. Therefore, this kind o f intergranular fracture ispseudo-intergranular fracture, for the fracture cracks donot nucleate and propagate along grain boundaries butin the PFZ.

C. Hydrogen Effect on Atomic Binding Energy

The atomic binding energy Eb can be cal led the ionicwork function,1241which is the work required to break upan atom from the specimen surface and remove it to aninfinitely distant a rea . According to the SIMS mecha-nism, sputtering and secondary ion emission o f an atomo f a certain element (such as A1 element) o f solid spec-imen can be represented as:

A1 (sol id)--A1+ (gas) + e - (reaction energy is Eb) (a)

A1 (sol id)--A1 (gas) (reaction energy is AH~) (b)

A1 (gas) - -A1+ (gas) + e (reaction energy is Ii) (c)

where AHs and 11 are the sublimation heat and the firstionization energy o f aluminum, respectively. Since (a)= (b) + (c),

Eb = ar ts + I1 [4]

According to energy conservation law and the mech-a n i s m o f SIMS analysis, 125,26,271 the energy o f primaryion Ep is the sum o f kinetic energy o f secondary ions EK,atomic binding energy Eb, and energy loss Eo:

Ep = EK + Eb + E0 [5]

When the primary beam intensity and other experimentalconditions are kept constant, the values o f Ep and E0 o fhydrogen-charged specimens are the same as those o funcharged ones. Therefore, the difference o f kinetic en-ergy o f secondary ion equals the difference o f atomicbinding energy o f hydrogen-charged specimens and hy-drogen uncharged ones; i .e . , AEb = - AEK. Hence , the

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Fig. 6--Fracture surfaces in high magnification (SEM): (a) uncharged, ~ = 3 .4 x 10-3/s; (b) charged, ~ = 3.4 x 10 J/s; (c) uncharged, g3.4 x 10-4/s; (d) charged, i = 3 .4 x 10 a/s ; (e) uncharged, g = 3 .4 x 10-S/s; and (f) charged, ~ = 3.4 x 10-~/s.

change o f kinetic energy o f secondary ions represents thechange o f atomic binding energy, and the increase o fkinetic energy o f secondary ions indicates the decreaseo f the binding energy. By taking the intersecting pointso f the curves on the retarding potential coordinate inFigure 7 as the retarding potential VR, AEbb = -AEK =- eAVR. In this experiment, AVR = -- 3.2 V; i.e. AEb= 3.2 eV. Similarly, the atomic binding energy differ-ence o f Li after hydrogen charging is obtained as AEb =2.3 eV. This means that the atomic binding energies o fmatrix aluminum and lithium decreased after hydrogencharging. To demonstrate the relative reduction o f atomicbinding energies o f A1 and Li after hydrogen charging,substituting AHs and 1~ with reported experimental re-sults, for simplicity, and taking the AHs and I~ o f pureAI element, i .e . , 3.23 eV~z81 and 5.98 eV[291 for the AHsand I1 o f the alloying element AI, respectively, then therelative reduction in atomic binding energy o f aluminumis

[AEb/(AHs + I,)] x 100 pct

= [3.2 eV/(3.23 eV + 5.98 eV)]

× 100 pct = 34.7 pct

Similarly, a f t e r substituting 1.59 eVtaSt and 5.39 eVt29~as AHs and 1~ o f li thium element , respectively, then,

[AEb/(AHs + I0] x 100 pct

= [2.3 eV/(1.59 eV + 5.39 eV)]

× 100 pct = 32.9 pct

Because li thium is an alloying element, the content o fwhich is less than 3 pct, and due to the shielding effecto f aluminum, the practical atomic binding energy re-duction o f li thiumaf te r hydrogen charging will surely bemore than the above calculated.

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Aluminium :~ 60 I6(3 - - Lithium ./

* ~ 5 0 - ~ -- 5 -

:30- 3o

g' - ' 2 0 - 0 c h a r ~ e d -~ ~ 20

01420 14~0 1500 15~ 1420 1460 1500 15/~)

Retarding p o t e n t i a l , V~ R e t a r d i n g p o t e n t i a l , VRFig. 7--Relation curves of secondary ion intensity of aluminum and Fig. 8--Relation curves of secondary ion intensity of lithium andretarding potential, retarding potential.

V. CONCLUSIONS

1. The tensile strengths (UTS, YS) o f hydrogen-charged8090 A1-Li alloy decrease linearly and the plasticities(RA,EL) of hydrogen-charged alloy drop greatly withthe decrease o f strain rate, but the plasticities are im-paired markedly af te r hydrogen charging.

2. Charged hydrogen increases the strengths o f 8090 A1-Lialloy slightly but markedly impairs the plasticity.

3. The susceptibility o f hydrogen embrittlement in-creases with the decrease o f strain ra te , and the sus-ceptibility o f the charged specimens is larger than thato f uncharged ones over the strain-rate range.

4. Secondary cracks on main fracture surfaces are de-lamination cracks along pancake-shaped grain bound-aries and perpendicular to the short transversedirection. There are more obvious delamination crackson the fracture surfaces o f hydrogen-charged speci-mens than on those o f uncharged ones.

5. The tensile fracture morphology o f charged speci-mens at slow strain rates (3.4 × 10-4/s and 3.4 x10 5/s) is G B B F , while that o f at fast strain rate (3.4X 1 0 - 3 / S ) is G B D F .

6. The atomic binding energies o f aluminum and lith-ium elements in the alloy decreased af te r hydrogen

charging, and the energy drop o f the former is greaterthan that o f the latter.

ACKNOWLEDGMENT

This research was supported by the National NaturalScience Foundation o f China .

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