Development of Porous Anodic Films on 2014-T4 Aluminium Alloy

Development of porous anodic films on2014-T4 aluminium alloy intetraborate electrolyte

A. Pakes a, G.E. Thompson a, P. Skeldon a,*, P.C. Morgan b

a Corrosion and Protection Centre, University of Manchester Institute of Science and Technology,

P.O. Box 88, Manchester, M60 1QD, UKb British Aerospace Ltd, Sowerby Research Centre, Filton, Bristol, BS12 7QW, UK

Received 10 June 2002; accepted 18 October 2002

Abstract

Anodic film growth on 2014-T4 aluminium alloy at 60 V in 50 g l�1 di-sodium tetraborate at

60 �C has been examined by transmission electron microscopy and Rutherford backscatteringspectroscopy. Initial film growth proceeds at relatively high efficiency on the initially etched

and desmutted alloy. During the subsequent period of current decline, the reactive electrolyte

species penetrate the outer film at preferred regions, establishing conditions for pore devel-

opment by field-assisted dissolution. In the alkaline electrolyte, such field-assisted dissolution

also appears to proceed locally, probably through mechanical disruption of the film, giving

rise to a feathered film morphology. The oxidation of copper from the alloy, in the presence of

an enriched layer of copper, developed largely by initial etching, also influences film mor-

phology through parallel oxygen gas generation, creating oxygen-filled voids. Such gas-filled

voids may rupture or be removed from the alumina film material through field-assisted dis-

solution at the pore base. In the former case, cracking allows access of the anodizing elec-

trolyte to the enriched alloy/film interface, with subsequent dissolution of the enriched layer

and local film growth; these give rise to lateral porosity in addition to that from pores passing

perpendicularly to the alloy surface. The efficiency of anodizing is about 12%, with losses from

Al3þ ion ejection, field-assisted dissolution, oxygen gas generation, film rupture, interface

dissolution and local film repair.

� 2002 Elsevier Science Ltd. All rights reserved.

Keywords: A. Aluminium; 2014 alloy; Anodizing; Anodic films

*Corresponding author. Tel.: +44-161-200-4872; fax: +44-161-200-4865.

E-mail address: [email protected] (P. Skeldon).

0010-938X/03/$ - see front matter � 2002 Elsevier Science Ltd. All rights reserved.

PII: S0010-938X(02 )00216-0

www.elsevier.com/locate/corsci

Corrosion Science 45 (2003) 1275–1287

mail to: [email protected]

1. Introduction

The 2000 series aluminium alloys are frequently favoured for aircraft structural

applications where damage tolerance and relatively high strength are paramount.

The strength of the alloy is achieved through heat treatment, which develops fine

strengthening precipitates of various structures and compositions, in addition to therelatively coarse equilibrium material formed through casting and thermomechanical

processing. Due to the differing electrochemical potentials of the matrix and pre-

cipitate materials, localized forms of corrosion are stimulated, rendering an inferior

corrosion resistance in comparison with relatively pure aluminium. In practice,

therefore, the 2000 series alloys may be clad and/or given surface treatments to

improve their corrosion resistance [1]. Traditionally, chromic acid anodizing has

been a preferred treatment in aerospace, providing protective anodic oxide films

that, where appropriate, may be sealed by hydrothermal treatment. Equally im-portant, anodizing generates an alloy surface that is suited to adhesive bonding, an

increasingly adopted method for joining aircraft components. Currently, there is

growing legislative pressure to develop alternative treatments to chromic acid ano-

dizing because of environmental problems associated with the chromate ion and the

costs of solution disposal. The task of finding replacements is complicated by the

unusual effectiveness of chromic acid, which allows formation of porous films of

appropriate morphology and provides corrosion inhibition to the substrate.

The choice, made by aircraft manufacturers, of pretreatments and anodizingtreatments for aluminium alloys partly reflects differing commercial and national

priorities and constraints. Thus, a range of specifications has emerged. In the search

for replacements for chromic acid anodizing, divergent approaches continue to be

evident. In Europe, there is interest in boric acid anodizing at 60 �C [2], while in theUS, a mixed sulphuric acid/boric acid process, carried out at lower temperature, has

been patented [3]. Other anodizing treatments, as well as alternatives to anodizing,

have also been proposed [4]. However, at the present time, the relative benefits of the

new treatments are unclear, especially with respect to practical applications whererequirements of corrosion resistance, adhesive bonding, wear resistance and fatigue

may be combined.

In the present study, the development of anodic films on 2014 alloy in alkaline

borate solution has been examined; transmission electron microscopy (TEM) has

been employed to reveal the development of the film and the underlying contribu-

tions to the relatively high population densities of flaws associated with the filmed

alloy substrate.

2. Experimental

2.1. Specimen preparation

Specimens, of dimensions 40� 10� 1 mm, were prepared from 2014 alloy sheet inthe T4 (solution treated and naturally aged) condition. The alloy composition is

1276 A. Pakes et al. / Corrosion Science 45 (2003) 1275–1287

given in Table 1. Following degreasing in acetone, the specimens were etched for480 s in 3 g l�1 sodium hydroxide solution at 25 �C, rinsed in deionized water, and thendesmutted in 50% (by vol.) nitric acid for 30 s at 25 �C, with final rinsing in deionizedwater and drying in a cool air stream. The etched and desmutted specimens were

anodized individually at a constant voltage of 60 V in stirred 50 g l�1 di-sodium

tetraborate solution, of pH 9.2, at 60 �C for selected times up to 1800 s, to developfilms with a range of thicknesses. During anodizing, the resultant current–time be-

haviour was recorded. The specimens were immediately removed from the electrolyte

at the termination of anodizing, rinsed in deionized water and finally dried in a coolair stream.

2.2. Specimen examination

Sections, about 15 nm thick, of the anodic films attached to the alloy substrate, were

prepared by ultramicrotomy and examined by TEM in a JEOL 2000 FX II instrumentwith energy dispersive X-ray (EDX) analysis facilities. The EDX analyses were carried

out using a probe of 20 nm diameter to optimize spatial resolution and count rate. Plan

views of the corresponding stripped films were examined after their removal from the

alloy substrate by amalgamation in mercuric chloride solution. Specimens were also

examined by Rutherford backscattering spectroscopy (RBS) using a 2.0MeV beam of

alpha particles supplied by the Van de Graaff accelerator of the University of Paris.

The beam current and diameter were 60 nA and 0.5 mm respectively, with the beam

incident normal to the specimen surface. The scattered particles were detected at 165�to the direction of the incident beam. The data were interpreted by the RUMP pro-

gram with scaling of the stopping power for oxygen by 0.88.

3. Results

3.1. Surface morphology after pretreatment

During etching of the 2014-T4 alloy in 3 g l�1 sodium hydroxide, copious gas

evolution was observed from the specimen surface. Upon completion of etching,

Table 1

Nominal composition of the 2014-T4 aluminium alloy

Element Weight%

Si 0.5–0.9

Fe 0.5

Cu 3.9–5.0

Mn 0.4–1.2

Mg 0.2–0.8

Cr 0.1

Ni 0.1

Zn 0.25

Al Balance

A. Pakes et al. / Corrosion Science 45 (2003) 1275–1287 1277

smut, of black appearance, was evident on the specimen. The subsequent dip in nitric

acid effectively removed the smut, to reveal a matt grey surface finish. A plan view of

the alloy surface following etching and desmutting reveals the topography of the

alloy surface comprising preferentially etched and relatively flat areas (Fig. 1(a)). Insection (Fig. 1(b)), the typical scalloped appearance of a preferentially etched area of

Fig. 1(a) is evident. Additionally, a continuous film, of thickness 5–10 nm, is present

over the etched alloy surface. A dark band, of thickness � 2 nm, located immediatelybeneath the film is associated with enrichment of heavy elements, primarily copper,

in the alloy.

3.2. Current density–time behaviour

The typical current density–time behaviour, with the current density expressed as

the average current density over the anodized surface, reveals an initial rapid surge

and subsequent decay of current density in the first few seconds, followed by a more

gradual rate of decay to a region of approximately constant current density (Fig. 2).

A small perturbation in the generally smooth decay, at about 100 s from the com-

mencement of anodizing, occurred for all specimens. The general form of the re-

sponse was reproducible between specimens, although currents differed by about10% at any particular time. The final current density, following anodizing for 1800 s,

was approximately 4 mA cm�2. Gas evolution was observed from the specimen

surfaces from almost the commencement of anodizing.

Fig. 1. Electron micrographs of the 2014-T4 alloy following etching in 3 g l�1 sodium hydroxide for 480 s,

and desmutting in 50% vol. nitric acid for 30 s: (a) scanning electron micrograph of the alloy surface;

(b) transmission electron micrograph of an ultramicrotomed section.


3.3. Transmission electron microscopy

The development of the anodic film over the macroscopic surface of the alloy is

revealed from observations of stripped films (Fig. 3). After 5 s, a time within the

region of current density decay after the initial surge, the relatively low magnification

transmission electron micrograph of Fig. 3(a) reveals film growth on the relativelyflat and scalloped regions of the originally etched substrate. Where the previous

regions merge, locally thicker film material is evident. The increased magnification

image of Fig. 3(b) shows the presence of approximately circular, light features,

particularly within the film developed over the scalloped areas of the substrate. Little

change in the film material is evident after 15 s (Fig. 3(c)) but, after 30 s, film growth

appears to be more uniform than at earlier anodizing times (Fig. 3(d)). The micro-

graph of Fig. 3(d) reveals the continued presence of approximately circular light

features, as well as evidence of fine porosity developing at the outer surface of theanodic film. Additionally, the film region displayed shows incorporation of fine

dispersoids, or associated regions of altered film morphology, within the stripped

film. With further anodizing for 120 and 240 s (Figs. 3(e) and (f)), similar appear-

ances are revealed, although the increased porosity at the outer surface of the film

and the associated regions of light appearance render difficult their distinction from

the approximately circular light features.

TEM of ultramicrotomed sections of the anodic film attached to the aluminium

alloy substrate has been employed to examine the various contributions to filmgrowth over the alloy surface (Fig. 4). Thus, after anodizing for 5 s, an approxi-

mately uniform film has developed over the undulating, etched alloy surface (Fig.

4(a)). Within the film, of thickness about 88 nm, light regions are present, of di-

mensions similar to the approximately circular features within the corresponding

Fig. 2. Current density–time behaviour during anodizing of the 2014-T4 alloy at 60 V in 50 g l�1 di-sodium

tetraborate at 60 �C.


stripped film. Local protrusions are evident at the film/electrolyte interface, located

above the light regions. A dark band, of width about 2 nm, similar to that present

Fig. 3. Transmission electron micrographs of stripped anodic films formed on the 2014-T4 alloy at 60 V in

50 g l�1 di-sodium tetraborate at 60 �C for different times: (a) 5 s; (b) increased magnification micrographof (a); (c) 15 s; (d) 30 s; (e) 120 s; (f) 240 s.


after pretreatment, is evident in the alloy immediately beneath the anodic film. Such

contrast, associated with increased electron scattering, indicates enrichment of rel-

atively heavy elements at such locations. Concerning the film/electrolyte interface,

other than for the presence of protuberances, there is little evidence of porosity. With

further anodizing for 15 s (Fig. 4(b)), the film thickness has increased to about 96 nm

and has a similar general appearance to the film of Fig. 4(a). However, at this an-

odizing time, there is distinct evidence for development of fine porosity at the film/electrolyte interface. The thickness of the outer porous region extends about 8–16 nm

from the film/electrolyte interface, with an underlying relatively compact film of

approximate thickness of 84 nm. The figure also reveals protuberances of film at the

film/electrolyte interface, between 10 and 80 nm in height. The protuberances are

Fig. 4. Transmission electron micrographs of ultramicrotomed sections of the anodic films formed

on 2014-T4 alloy at 60 V in 50 g l�1 di-sodium tetraborate at 60 �C for different times: (a) 5 s; (b) 15 s;(c) 240 s; (d) 360 s; (e) 1200 s.


generally associated with voids in the film, which suggests they are probably formed

as a result of the pressure of oxygen gas within the voids being sufficient to deform

and then rupture the anodic film. After 240 s, the inner compact film has a thickness

of about 86 nm, but the outer porous region has clearly thickened to approximately

100 nm (Fig. 4(c)). As indicated previously, light features are present within the film

material, and the dark band, enriched in alloying elements, is evident in the alloyimmediately beneath the anodic film. After anodizing for 360 s, the inner compact

film material remains of similar thickness to that at earlier anodizing times, with

further thickening of the porous region to 150 nm (Fig. 4(d)). With regard to the

appearance of the film material, this is again similar to that at the previous anodizing

times other than for the presence of light features passing linearly across the film

section, approximately mid-way between the film/electrolyte and alloy/film inter-

faces. With anodizing for 1200 s the overall film thickness has generally increased,

although the inner region of compact film has maintained an approximately constantthickness of about 90 nm. With continued anodizing to 1800 s (Fig. 4(e)), the total

film thickness has increased to about 820 nm, with a barrier layer thickness of ap-

proximately 80 nm. However, in this region of the section, the alloy/film interface

appears highly irregular, with the anodic film penetrating further into the alloy in

preferred regions. Fig. 5 reveals the non-uniformity of film growth over the mac-

roscopic alloy surface after anodizing for 1800 s. It is evident that the anodic film has

penetrated relatively deeply into the substrate at preferred regions, possibly related

to the initial lightly etched regions of Fig. 1.

Fig. 5. Transmission electron micrographs of an ultramicrotomed section of the 2014-T4 alloy after

anodizing for 1800 s at 60 V in 50 g l�1 di-sodium tetraborate at 60 �C, revealing non-uniform film

development over the alloy surface.


3.4. Rutherford backscattering spectroscopy

The films examined by TEM were also probed by RBS. The interpretation of the

RBS spectra assumed that the films were composed primarily of alumina of density

3.1 g cm�3, corresponding to that of non-porous anodic alumina. The results reveal

that, for anodizing times up to 60 s, the film thicknesses determined using ultrami-crotomed sections and RBS data are similar; this indicates that porosity is relatively

low (Fig. 6 and Table 2). For anodizing times greater than 60 s, the film thickness

determined from the ultramicrotomed sections exceeds that determined from RBS by

roughly 5–15%, which is attributed to the developing porosity revealed in ultrami-

crotomed sections.

Fig. 7(a) displays a typical spectrum for a film formed to 60 s, revealing the

presence of alloying element species in the anodic film. Due to the closeness of the

atomic masses of chromium, manganese, nickel, iron, zinc and copper, it is notpossible to give an unequivocal interpretation of the data. However, the data suggest

Fig. 6. Variation of film thickness, determined using TEM of ultramicrotomed sections and RBS, with

time, for anodizing 2014-T4 alloy at 60 V in 50 g l�1 di-sodium tetraborate at 60 �C.

Table 2

Summary of film thicknesses for anodic films formed on the 2014-T4 aluminium alloy at 60 V in 50 g l�1 di-

sodium tetraborate at 60 �C for various times.

Anodizing time (s) Film thickness from TEM (nm) Film thickness from RBS (nm)

5 88 105

15 96 110

60 128 125

120 176 135

240 187 170

360 238 220

600 280 265


strongly that copper, the heaviest and most abundant of the alloying elements, is

present in the film since a step can be resolved in the spectra which can only be

assigned realistically to this copper. The alternative of zinc is disregarded due to its

low concentration in the alloy, although zinc is expected to make a relatively small

contribution to the yield in this region of the spectra. The atomic ratio of Cu to Al

species in the film is about 9� 10�3, which is about one-half that in the alloy. Thislevel of reduction in copper content between the film and the alloy is typical of that

found for formation of barrier-type films on dilute aluminium–copper alloys, and isassociated with the faster outward migration of copper species across the barrier

layer relative to that of Al3þ ions, which leads to loss of the copper species to the

borate electrolyte on reaching the film/electrolyte interface [5].

In spectra of thicker films, i.e. those formed for 360 s (Fig. 7(b)) and beyond, the

yields reveal a step suggesting an increased concentration of copper species in the

inner part of the film. However, contributions from iron and other alloying element

species may also arise. The step height, if interpreted as due solely to iron species, the

second most abundant alloying element, indicates an Fe:Al ratio of 2:2� 10�2,which is greater than the ratio for the alloy. EDX analyses of the anodic films (not

shown) revealed mainly copper species in the films, indicating the step is mainly due

to their increased concentration in the inner film layers.

Notably, the RBS spectra do not reveal evidence of the presence of oxygen gas

within the film, a feature of barrier film growth on dilute aluminium–copper binary

alloys. However, this can be explained by the lack of retention of such generated gas

due to rupture of the film material as a result of significant pressures in the gas-filled

voids, assisted possibly by the development of porosity in the film. The initial releaseof oxygen gas in the period of transformation from initial non-uniform growth to

steady-state anodizing is considered also to influence the early region of the current

density–time behaviour of Fig. 2.

Finally, all spectra reveal a peak associated with an enriched alloy layer adjacent

to the alloy/film interface. The precise interpretation of the peak is complicated due

Fig. 7. Experimental (points) and simulated (line) RBS spectra for the 2014-T4 alloy after anodizing at

60 V in 50 g l�1 di-sodium tetraborate at 60 �C for different times: (a) 60 s; (b) 360 s.


to possible influences of iron and other elements of similar mass, which may be

present in the anodic film. However, a good fit can be obtained for an assumed

enriched alloy layer of 2 nm thickness containing about 35 at.% Cu. This level of

enrichment is similar to that found previously for binary aluminium–copper alloys of

similar copper content to that of the present alloy [5].

4. Discussion

4.1. Influences of alloy thermomechanical processing and pretreatment

After caustic etching and desmutting, the surface reveals locally rough regions, of

scalloped appearance, and adjacent, comparatively flat regions. Such regions are

thought to arise through alloy fabrication and associated heat treatment; the local

rough regions result from relatively rapid etching behaviour compared with the light

regions. The influence of thermomechanical processing on the surface of aluminium

alloys has been considered recently [6–8], with deformed surface layers developedlargely during hot rolling; such layers deformed may be removed by appropriate

cleaning. However, unlike the 2014 alloy here, deformed layers were present gen-

erally over the alloy surface. Further, through alkaline etching, solid-solution al-

loying elements, principally copper, enrich in the alloy beneath the alumina film. The

desmut in nitric acid removes copper from the film surface, but the enriched layer

remains.

4.2. Anodizing behaviour

Previous investigations of the anodizing of superpure aluminium under similar

conditions revealed that the initial film growth proceeds at high efficiency [9]. Sub-

sequent penetration of the outer film regions by the electrolyte results in concen-

tration of the field at preferred sites and pore initiation. The resultant porous films

revealed a feathered morphology, similar to that found after anodizing in chromic

acid, due to the intense field at the pore base mechanically disrupting the film. For

the alloy situation, broadly similar features are revealed, but the thermomechanical

processing of the alloy and its surface pretreatment have an influence as well as thealloy composition.

4.3. Anodic film growth

The enriched alloy substrate, with the residual film developed on the etched and

desmutted surface, undergoes anodic polarisation by anodizing. During the initial

current surge, growth of barrier-type anodic oxide proceeds at relatively high effi-

ciency, with film material forming at the film/electrolyte and alloy/film interfaces byAl3þ egress and O2� ingress under the high field. Additionally, in the presence of the

enriched layer of copper, copper may enrich further prior to achieving the critical

concentration at which it oxidizes [10]. RBS reveals copper species are present in the


film after relatively early anodizing times, suggesting that the level of enrichment

after etching is close to the critical concentration. From the magnitude of the Cu2þ–

O bond energy relative to that of the Al3þ–O bond, copper species are outwardly

mobile at a higher rate than aluminium ions [11]. This leads to ejection of copper

species at the film/electrolyte interface in the borate electrolyte.

After the initial film growth, from the period of current density decay and beyond,when the field strength across the alumina film decreases, the outer film surface is

penetrated by electrolyte species. This effectively thins the alumina film in preferred

regions, allowing the field strength to increase locally, with associated current flow.

At preferred regions, with locally high field strengths as a result of the local film

geometry, field-assisted dissolution of the alumina is promoted. In alkaline borate

electrolytes, where the reactivity of the electrolyte species to the alumina appears less

than in the common acidic electrolytes, this contributes to the feathered pore mor-

phology.

4.4. Oxygen gas generation

In the presence of the enriched alloy layer, copper is oxidized to form Cu–O units,

which, through ionic transport, are dispersed across the amorphous alumina film

material. Oxidation of copper is also associated with oxidation of O2� ions within thefilm, giving rise to oxygen gas-filled voids. Such gas-filled voids influence ionic

transport locally, effectively increasing film resistivity, and contributing to the de-

tection of an increased copper ion content in the inner film regions. Further, gas

builds up in the voids to extremely high pressure, i.e. about 200 MPa [10], with

eventual film rupture. However, in the situation where a porous anodic film devel-

ops, the barrier layer film material undergoes field-assisted dissolution at the pore

base which is balanced by growth at the alloy/film interface. Consequently, when the

pore base reaches the voids, the gas is released with a local alteration of the poremorphology. In instances where the gas pressure builds up to high levels, film rup-

ture and cracking may occur, with the electrolyte accessing the alloy surface at high

potential. Here, the enriched interfacial region undergoes dissolution and new film

then develops above the underlying alloy. This gives rise to enhanced porosity in the

film, with this porosity orientated parallel to the alloy/film interface, unlike the true

pores which pass perpendicularly to the alloy surface.

4.5. Efficiency of film growth

The film thickness achieved after anodizing for 1800 s, i.e. 800 nm, compares with

an expected thickness of 4000 nm for anodizing at 100% efficiency at a constant

current density of 4 mA cm�2. Thus, the average efficiency of anodizing of the alloy

over 1800 s is about 12%. The losses arise from Al3þ ion ejection and field-assisted

dissolution at the pore base, necessary conditions for porous film formation. Ad-ditional losses result from copper oxidation and ejection at the film/electrolyte in-

terface, but the major contribution arises from the parallel processes of oxygen

generation, film cracking, interfacial dissolution and local regrowth.


5. Conclusions

1. Anodizing 2014-T4 etched aluminium alloy at 60 V in 50 g l�1 di-sodium tetra-

borate at 60 �C results in the development of a porous anodic alumina film atlow efficiency. The porous film morphology arises from field-assisted dissolution

at the pore base in the relatively mild electrolyte and film growth at the alloy/filminterface.

2. Oxygen generation is associated with the development of an enriched alloy layer,

containing mainly copper, which remains immediately beneath the anodic film

during anodizing; gas generation proceeds in parallel with oxidation of copper

and incorporation of outwardly mobile copper ions into the anodic alumina.

3. The oxygen gas bubbles influence the porous film morphology. The most signifi-

cant influence is where such regions rupture and the electrolyte has direct access to

the alloy surface at high potential. This leads to dissolution of the enriched regionand subsequent growth of new film material, leading to regions of lateral porosity,

in addition to the feathered pores that pass perpendicularly to the alloy surface.

Acknowledgements

The authors are grateful to the Engineering and Physical Sciences Research

Council (UK) for the award of a studentship to A. Pakes and British Aerospace Ltd,

Sowerby Research Centre, for support of the research.

References

[1] S. Wernick, R. Pinner, P.G. Sheasby, The Surface Treatment and Finishing of Aluminium and its

Alloys, Finishing Publications Limited, Teddington, 1996.

[2] D. Marchandise, I. Belliot, R. Marrugat, Aerospatiale, French Patent 9517536, 22 Dec 1995.

[3] C.-M. Wong, Y. Moji, Boeing Corp, US Patent 4894127, 24 May 1989.

[4] C.J.E. Smith, K.R. Baldwin, S.A. Garette, M.C. Gibson, M.A.H. Hewins, P.L. Lane, Proceedings of

the 1st International Symposium on Aluminium Surface Science and Technology, Antwerp, Belgium,

1997, p. 266.

[5] H. Habazaki, M.A. Paez, K. Shimizu, P. Skeldon, G.E. Thompson, G.C. Wood, X. Zhou, Corros.

Sci. 38 (1996) 1033.

[6] M. Fishkis, J.C. Lin, Wear 206 (1997) 156.

[7] H. Leth-Olsen, J.H. Nordlien, K. Nisancioglu, J. Electrochem. Soc. 144 (1997) L197.

[8] G.M. Scamans, A. Asfeth, G.E. Thompson, X. Zhou, Proceedings of the 2nd International

Symposium on Aluminium Surface Science and Technology, Manchester, UK, 2000, p. 9.

[9] A. Pakes, G.E. Thompson, P. Skeldon, P.C. Morgan, K. Shimizu, Trans. Inst. Met. Finish 77 (1999)

171.

[10] P. Skeldon, G.E. Thompson, G.C. Wood, X. Zhou, H. Habazaki, K. Shimizu, Phil. Mag. A 76 (1997)

729.

[11] H. Habazaki, X. Zhou, K. Shimizu, P. Skeldon, G.E. Thompson, G.C. Wood, Electrochim. Acta 42

(1997) 2627.


Development of Porous Anodic Films on 2014-T4 Aluminium Alloy

Documents

Transcript of Development of Porous Anodic Films on 2014-T4 Aluminium Alloy