JJW o

NBSiR 80-2142 (Aluminum Association)

D. S. Lashmore

Corrosion and Electrodeposition Group

Chemical Stability and Corrosion Division

National Measurement Laboratory

U.S. Department of Commence

National Bureau of Standards

Washington, DC 20234

November 1980

Final Report

Prepared for The Aiuminum Association

818 Connecticut Avenue, N.W.

Washington, DC 20006

on Aluminum

AUG 5 1981

The American Electropiaters Society

1201 Louisiana Ave.

Winter Park, FL 32789

NBSIR 80-2142 (Aluminum Association)

PLATING ON ALUMINUM

D. S. Lashmore

Corrosion and Electrodeposition Group

Chemical Stability and Corrosion Division

National Measurement Laboratory

U.S. Department of Commerce

National Bureau of Standards

Washington, DC 20234

November 1980

Final Report

Prepared for The Aluminum Association The American Electroplaters Society

818 Connecticut Avenue, N.W. 1201 Louisiana Ave.

Washington, DC 20006 Winter Park, FL 32789

U.S. DEPARTMENT OF COMMERCE, Philip M. Klutznick, Secretary

Jordan J. Baruch, Assistant Secretary for Productivity, Technoiogy, and Innovation

NATIONAL BUREAU OF STANDARDS, Ernest Ambler, Director

Table of Contents

Page

1. Introduction 1

2. Program Summary 2

3. Phosphoric Acid Anodizing Prior to Plating 10

3.1. Adhesion Data 10

3.2. Microstructure of the Anodic Film 17

3.2.1. Microstructure of nickel plated onto the anodic 19

film

3.2.2. Initial Stages of Oxidation 25

3.3. Electrochemical Measurements 32

3.3.1. Current Density versus Applied Potential 32

3.3.2. Scanning Voltammetry 33

4. Immersion Deposition 37

4.1. Scanning Voltammetry 37

4.2. Dissolution Data 41

4.3. Discussion 44

5. Adhesion Theory 45

5.1. Adhesion of Metallic Coatings 45

5.2. Pre-existing Crack--Griffith Criterion 52

5.3. Psuedomorphism: the critical radius and the critical 56

thickness

List of Figures 61

References 36, 60
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- 1 -

1. Introduction:

This report covers the period from about July 1979 to February

1980 and includes, not only a discussion of work on both anodizing

and immersion deposition, but also a summary of the results obtained

over the last three years. A general theory of adhesion of metal

coatings on metal substrates has been developed and adhesive energies

were calculated for epitaxial zinc on aluminum. The case of a

preexisting crack in the interface was considered.

The work on anodizing has resulted in a greatly improved

understanding of adhesion mechanisms. The morphology of the initial

stages of anodic film formation on aluminum alloy 7146-T6 has been

characterized, yielding results rather different than predicted

from commonly accepted theories. The morphology of nickel deposited

into the pores of the anodic film has been characterized for the

first time, it is believed, by di rect TEM studies of cross sections

of bulk specimens on which adhesion measurements have been made.
- 2 -

2. Program Summary:

The reader is urged to refer to particular progress reports

and to Plating 37 (1980) and the Journal of the Electrochemical

Society 1 27 , 573 (1980) for more complete information on the first

two years' results. During the three years that the Aluminum

Association project Decorative Chrome Plating of Aluminum has been

carried out, three major areas have been worked on: immersion

deposition, anodizing, and the theory of adhesion between metal

coatings and substrates. A summary of the results follows:

2.1. Immersion Deposition:

2.1.1. Substrate Orientation :

It was shown in Progress Reports (P. R.) II-IV that the

substrate crystal structure plays an important role in both

the zincate and stannate processes. The dissolution of the

aluminum in the caustic zincate solutions is a sensitive

function of the crystal 1 ographic orientation. Moreover, zinc

was shown to form epitaxial deposits on all aluminum principal

planes. The bronze in the stannate process was shown to

exhibit a strong preferred orientation.

2.1.2. Cleaning Process :

In virtually every process for immersion deposition, a

caustic cleaning step is specified. The highly deformed

surface usually present on most material interferes with

epitaxial phenomena. In addition, this highly deformed layer

would be expected to dissolve more rapidly than less deformed

material, permitting a thick zinc deposit to form preferentially at regions of high deformation. The etch cleaning

step not only removes the highly deformed material but also
-3-

attacks the same families of planes as the zincate and stannate

electrolytes. This step is a necessity, if the balance between

aluminum dissolution and uniform zinc deposition is to be

mai ntai ned

.

2.1.3. Alloy Constituents :

It was shown that even small traces of magnesium determine

the sites at which large zinc crystallites grow. The epitaxial

phenomena observed on pure aluminum were also observed on one

alloy (6061-T6) and probably exist on others (P. R. II & III).

Investigations of adhesion correlations with alloying constituent

were discussed in P. R. IV. Pure aluminum exhibited greater

adhesion than did the alloys investigated here.

2.1.4. Effect of Ferric Chloride and Rochelle Salts :

It was shown that small additions of ferric chloride and

Rochelle salts have a major effect on the deposition process.

When absent, it was shown that three atoms of zinc are deposited

for every two atoms of aluminum dissolved. When present, it

was shown that almost nine atoms of aluminum are dissolved for

each atom of zinc deposited. (See P. R. II.) It was shown

that iron is deposited (by immersion or perhaps chemisorption)

before the bulk of the zinc. Following a nitric acid immersion

(the double zincate procedure), it was shown by Rutherford

back scattering experiments (P. R. Ill) that about two monolayers

of iron and two monolayers of zinc were left on the surface.

Transmission electron microscopy of double zincated specimens

indicated the existence of an iron zinc alloy formed just at

the aluminum surface. (See P. R. I-IV.)
-4-

2.1.5. Morphol oqy :

The morphology of immersion deposits was characterized by

optical, SEil and TEM techniques, both on pure aluminum and on

a number of alloys. See P. R. I-IV, Plating 37 (1980),

and J. Electrochem. Soc. 1 27 , 573 (1980). Zinc forms epitaxial deposits on the aluminum substrates so that Zn

(0001 ) < 1010 > //Al (100) < 010 >, Zn:

(0001 ) < 1120 > //Al (111) < 110 > and Zn :

(0001 ) < 1120 > //Al (110) < 111 >

2.1.6. The Stannate Process :

It was shown by TEM that tin crystallites cover the

surface following the stannate immersion process. The tin was

shown to deposit in three distinct morphologies (See P. R.

O II.): (1) the formation of a fine grained deposit (<20A)

completely covering the surface; (2) at a later time, the

formation of small crystallites in discrete areas; and, finally,

(3) the growth of large, almost spherical clusters. It seems

that stage III is just barely reached after about 45 seconds

in the electrolyte.

The bronze layer deposited as part of the

stannate process was shown to exhibit a strong preferred

orientation. However, extensive work on this phenomenon was

not done because of the proprietary nature, of the commercial

process of which it is a part.

Significantly, it was shown that there are important

adhesion variations from the top to the bottom of vertically
-5-

plated specimens. The greatest adhesion occurred on the lower

part. Apparently, this phenomenon was also revealed in the

field by corrosion variations on plated bumpers.

Adhesion on a number of alloys was characterized by peel

testing. Deposition on single crystal spheres indicated that

a similar deposition mechanism existed with the zincate process,

(P. R. II, III & IV).

2.1.7. Fi el d Ion Microscopy ( F . I . M .

)

Techniques were developed to image zinc coated aluminum

in the F.I.M. (see P. R. IV). The zinc was observed to form

blade shaped deposits on the aluminum tips, probably due to

selective deposition on certain crystal 1 ographic planes.

2.1.9. Scanning Voltammetry and Acid Dissolution Data :

Scanning voltammetric data indicated that the deposition

of an iron-zinc or iron layer precedes the deposition of the

zinc. By monitoring the substrate potential as a function of

time, a modified procedure for depositing a thin layer of zinc

or zinc-iron alloy was developed (see Interim Report, Aug.

1 979, and section 4.2 of this report)

.

2.2. Anodizing Pretreatments

2.2.1. Effect of Structure

A procedure was developed to enable direct TEM micrographs

of cross sections of anodized aluminum coated with nickel.

Apparently, this is the first time that specimens of this type

have been prepared, though it should be noted that there has

been much published on replicas of anodic films and on
- 6

-

microtomed anodized aluminum. The data shown in this report

reveals n£ penetration of the barrier layer by the nickel,

ruling out metallic bonding as an adhesion mechanism.

The initial states of pore formation were characterized

by direct TEM studies of anodized 7021 -T63 and 7146-T6. The

initial stages of growth were shown to be much different from

subsequent stages. The anodic oxide seemed to form small

square regions which grew together like building blocks on a

(100) aluminum substrate (see section 3.2.2 of this report).

This evidence indicates the important role played by the

structure of the substrate during the initial stages of growth.

The classical model of pore growth requires a field assisted

dissolution of the aluminum at places where the oxide is

thinnest, and such a model is supported by the results obtained

during this study. However, in the initial stages of oxide

formation the thin areas are not necessarily associated with

pore formation; indeed, it was shown that oxide growth can be

quite remote from the dissolution sites, as can the initial

pore structures

2.2.2. Effect of Anodizing Parameters on Adhesion :

It was shown that adhesion is an approximately linear

function of the applied voltage. Since the cell diameter also

increases in a linear manner with the applied voltage, a

relationship was thought to exist between adhesion and pore

diameter. Moreover, the greater the cell diameter, the less

oxide there is per unit area of the coating. This implies

that the strength of the oxide itself is the determining
-7-

factor (or weak link) in adhesion of subsequent coatings to

the anodic film. It should be noted that the pore structure,

especially at higher voltages, does not vary much as one

changes from one of the three major electrolytes used for this

process (sulfuric acid, oxalic acid, and phosphoric acid) to

another even though adhesion of subsequent coatings can vary

by more than an order of magnitude between the various el ectrolytes

.

Therefore, the adhesion does not appear to be dependent upon

the pore structure, but, rather, upon the nature of the oxide

itself. The composition of the anodic film, varies considerably

with the alloy and with the electrolyte.

It was shown (see section 3 .1 of this report) that the

^

the use of an 0.5 percent solution of HF in place of sulfuric

acid on 7146-T6, results in increased adhesion.

2.3. Theory of Adhesion for Metallic Coatings on Metallic Substrates:

2.3.1. A theory of the adhesive energy for metalliccoatings on metallic substrates was discussed om detail (see

section 5 of this report). The theory takes into account the

effect of structure through a parameter 'a' which represents

the lattice parameter i_n the interface. This parameter is

determined by the orientation of both the substrate crystal

and the coating crystal . The adhesive energy as a function

only of the structure is

E ad

K G a^

0

(2TT)"d

[

1 ]
- 8 -

For f.c.c. materials, K = 3 , is the shear modulus in the

interface which can be determined from the binding energy

between the two metals and d is the distance of closest approach. The lattice parameter in the interface, a, is given

by'.

a

2a a

c s a + a

c s

[ 2 ]

2.3.2. In many cases perfect epitaxy is not possible; therefore,

the misfit between the coating and the substrate has to be

taken into account. It was shown (see section 5.1 of this

report), that the adhesive energy in the case of significant

misfit is given by

9 KG a

E .4 = E . - E (2ir)^d

^ [3]

where is the dislocation energy per unit area and E^ is

the distributed strain energy per unit area. A special case

was calculated for zinc coatings on aluminum whose epitaxial

relationships were characterized during this program.

2.3.4. The case of a preexisting crack of length 2 c, in the

interface was considered, and a Griffith type equation for the

minimum stress, '^critical, necessary to propagate this crack

(and thus to cause adhesive failure) was derived as^
- 9

-

This equation will be improved in the future to include

the effect of strain energy which may be in many cases an

important term.
- 10 -

III. Phosphoric Acid Anodizing Prior to Plating;

3.1. Adhesion Data

A number of experiments concerned with anodizing of

aluminum alloys were performed. A great deal of data has been made

available by Uittrock, so the experiments described here were

designed to compliment the available literature. The experiments

were specifically set up: (a) to characterize the effect on adhesion

of variations in the applied anodizing voltage. Surprisingly, no

direct data on this subject seems to be available, (b) to examine

the effect of post-anodizing treatment on adhesion, and (c) to

extend Wittrock's measurements of adhesion as a function of anodizing

electrolyte temperature.

3.1.1. Adhesion versus anodizing voltage!

The alloy chosen for this study was 7146-T6 because it

is in wide use as a bumper alloy. Studies on 7129-T63 will be

reported in the future. Specimens of 5 cm x 10 cm x 0.64 cm

were buffed, and degreased with trichl oroethl eyne in an

ultrsonic cleaner. Following this treatment, the specimens

were further polished with 0.01 ym MgO by hand, given a 30

second treatment in 50 percent nitric acid in an ultrasonic

cleaner, rinsed in distilled water, and immediately anodized.

The total anodizing charge (for adhesion test specimens) was

2 standardized to about 20 coulombs/cm regardless of anodizing

time. An amp-hour meter was used to monitor the total charge.

Following the anodization process, the specimens were given a
- 11 -

post treatment, typically in an electrolyte of 2 g/L sulfamic

acid held at 55-60 °C. The specimens were then rinsed in

distilled water and placed in a nickel sulfamate electrolyte,

maintained at a pH of 2.5 for 60 seconds prior to applying a

potential. This potential was increased at a very slow rate

in such a way that the current density went from 0 to 4.3

2 amps/dm in about 5 minutes. Adhesion was measured by peel

testing. All specimens were plated with a 150 ym thickness of

nickel

.

The adhesion was found to be an approximate 1 inear function

of anodizing voltage (figure 1). The extended curves intercept

the x-axis at about 15 volts, indicating that 15V is the

threshold below which the adhesion is zero. The slope of the

curve, however, changed with the concentration of the acid.

The current density versus applied voltage has been superimposed

on this data, showing that the greatest adhesion occurs when

the current density is about to become unstable. The electrolyte

was agitated mechanically. Because the limiting current

density is dependent on agitation, the curves shown in figure 1

may not represent the maximum obtainable current density, even

though the mixing was quite virorous.

The linear increase of adhesion with anodizing voltage is

unexpected if the major bonding forces are Van Der Waals, as

thought by Alexsandrov (1), or even if they are due to ionic

or covalent bonding. The cell diameter as a function of

applied voltage, figure 2 as calculated from the data of

Keller, Hunter and Robinson (2). It should be noted that
- 12 -

Figure caption for figure 1 (page 13)

Figure 1. Adhesion and current density as a function of anodizing

voltage (a) 10 percent ohosphoric acid, 60 ®C, (b) three percent phosphoric

acid, 60 ®C, and (c) 10 percent phosphoric acid, 30 ®C.
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- 14-

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-15-

O'Sullivan and Wood (3) have shown that the pore diameter is

also linearily dependent on applied voltage. If the bonding

between the pore walls were solely responsible for adhesion,

then, as the area available for bonding increased with decreasing

voltage, the adhesion would also be expected to increase. The

data in figure 2, compared to figure 1 , show exactly the

opposite trend. As the area for bonding increases, the

adhesion decreases. The adhesion as a function of pore diameter

is shown in figure 3. The linear dependence of adhesion on

cell size shown in figure 2, implies that the oxide is the

weak link. As the proportion of oxide in the anodic film

decreases at higher voltages, the adhesion increases. This

finding is consistent with other evidence. For example, the

microstructure of the anodic film as shown by many investigators

is essentially the same on all alloys, even on pure aluminum.

At higher voltages, as shown by Keller, Hunter and Robinson,

the pore microstructure does not vary much between different

anodizing electrolytes. For example, oxalic acid, sulfuric

acid and phosphoric acid all give rise to similar pore diameters

at voltages of around 60 volts, though adhesion variations

between alloys and even electrolytes can be more than three

orders of magnitude. This new evidence indicates that efforts

to improve adhesion of the anodic film should be directed at

increasing the strength of the oxide itself. It should be

noted that metallic bonding through the base of the pore by

the nickel has been ruled out by the TEM evidence.
- 16 -

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300

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-17-

3.1.2. The Effect of Post Anodizing Treatment:

Extensive experiments on post anodizing treatments have

not yet been done. However, preliminary experiments indicate

that dilute HF (0.5 vol . percent) solutions provide greater

adhesion than the commonly used sulfamic acid. However if

more concentrated solutions are used, say 2 g/L, then the

adhesion drops off. If the concentration of the sulfamic acid

is increased to 4 g/L, then adhesion increases by about 25

percent. The strong HF solutions obviously attacks the oxide,

more rapidly than the sulfamic acid.

A possible explanation for the effect of the post anodizing

treatment is that the acid removes a weak boundary layer from

the pore walls and further cleans out those pores filled with

reaction products.

3.2. Microstructure of the Anodic Film:

In order to determine the mechanism of adhesion of nickel,

plated in a porous anodic film, a transmission electron microscopy

(TEM) investigation of the microstructure was undertaken.

Specimens of several different kinds of alloys were investigated

and were prepared in either of two ways,: (a) so that the

interface was perpendicul ar to the electron beam, or (b) so

that the interface was parallel to the beam. Data from case

(b) have not been reported in the literature. Therefore, this

investigation represents, the first time that a direct TEM

study of nickel plated into the anodic film has been characterized

by TEM. The few TEM studies of anodized specimens (not nickel

plated) prepared with the interface parallel to the beam were
-18-

either of replicas or were of specimens prepared by ultramicrotome techniques.

Anodic film on both 7129-T63 and 7146-T6 were investigated.

The specimens were anodized according to Wittrock's procedure.

The starting geometry was 10 cm x 10 cm x 0.64 cm (4" x 4" x

1/4") for 7146-T6 and about 5 cm x 1 5 cm x 0.64 cm (2" x 6" x

1/4") for 7129-T63. The large specimens were used for adhesion

determination prior to further preparation. It was thought

important to know how an interfacial structure under study was

related to a given adhesion value. Following adhesion testing,

an untested section near the center of the panel was nickel

plated for an extended time in the same sulfamite nickel

electrolyte used for peel testing. The plating was continued

until 5 mm of nickel was built up.

Foils were then prepared from this specimen so that the

interface was perpendicular to the surface — the interface

would appear as a line through the foil. Two techniques were

used to cut foils: (a) a spark cutter, and (b) a LEGO - VARI

CUT (VC-50) low speed diamond cut off saw. Foil thickness

varied from 75 to 250 um. The foils prepared by spark cutting

were further polished using standard metal lographic techniques,

with the final polish being 0.01 ym MgO. The diamond cut-off

specimens were used without further processing. Both sets of

foils were then thinned to electron transparency in an ion -

milling machine with 3.5 to 4 KeV Argon ions. This latter

technique required milling times in excess of a week, operating

24 hours/day. The thinned foils were then examined in a
- 19 -

Phillips EM 200 A at 100 keV, equipped with the usual anticontamination stage and a double tilt specimen holder.

3.2.1. A representative micrograph of 7129-T63 anodized in 3

percent phosphoric acid is shown in figure 4. The adhesion,

measured by peel testing, exceeded 75 M/m (100 #/in.) The

O average cell size, at the interface , is about 1250 A, yielding

9 2

a cell density of 6.4 x 10 /cm . The tip radius is about 850

OA. Several important structural features of the anodic oxide -

nickel system are revealed here. First, the nickel has penetrated

the pores completely and, secondly, there appears to be a

continuous oxide barrier between the nickel and the aluminum.

O The nickel grain size in this specimen is about 400 - 500 A.

For this micrograph, the substrate was not oriented for

dislocation contrast, so the effect of the al umi num ' s . structure

on the growth of the anodic film was not apparent.

The aluminum's structure has a very important effect on

the growth of the anodic film, as shown in figure 5. This

micrograph was exposed and printed to bring out the dislocation

structure. It is seen that, in a number of places, the oxide

o grows along dislocations. Regions of oxides of 1 50 A diameter

appear to precede the aluminum/oxide interface by as much as

o 1200 A. Moreover, many of the oxide regions appear to have

four fold symmetry (though no diffraction from them was observed).

As the ion thinning proceeded into the nickel, it became

apparent that the pores (filled with nickel) were not at all

uniform, but underwent a great deal of branching, as shown in

figure 6 abc. This branching is probably caused by a voltage
^ irtpl

s:

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- 20-

1 1 10 0 nm

Figure 4. A transmission electron micrograph of 7129 T63 anodized in

3% phosphoric acid then plated with nickel from a sulfamate electrolyte.

- 21 -

Figure 5. A TEM micrograph of 7129 T63 showing the growth of the anodic

oxide along dislocations in advance of the barrier layer. The dark area coeresponds to nickel plated into the porous oxide.

- 22 -

Figure 6a. A low magnification bright field TEM micrograph of nickel

plated into a porous anodic oxide exhibiting typical pore branching.

- 23 -

Figure 6b. A TEM bright field micrograph of the same srea as shown in

figure 6a but at higher magnification.

- 24 -Hi

Figure 6c. A Dark field micrograph of the same area as shown in 6a taken

with the objective aperature over the diffuse oxide ring.

-25-

transient at the initial stages of film growth. Certainly

this type of microstructure would result in a strong mechanical

bondi ng

.

3.2.2. The initial stages of growth of the anodic film on specimens

of 7146-T6 were studied by TEM, with the beam perpendicul ar to

the interface. The mechanism of initial film formation is

apparently different from growth at a latter stage. Several

surprising results were observed in figure 7. A large number

of rectangular and even square regions of oxide are observed

on the (100) plane of 7146-T6. This observation indicates

that a strong role is played by the structure of the substrate.

These regions of oxide grow in such a way as to form pores, as

indicated on the micrograph. Even the large pores show evidence

of being made up of rectangular 'building blocks' of oxide.

The second major feature observed during the initial stage of

film formation is a large number of fringes which may be Moire

fringes. If subsequent study shows this to be the case, then

it is clear that the oxide forms, at least at the initial

stage, an epitaxial (or perhaps pseudomorphic) layer before

becoming amorphous. No evidence of crystalline oxide was

detected in the diffraction patterns.

The white areas shown in figure 7 are etch pits. The

classical model of pore growth requires field assisted dissolution

at places where the oxide is weakest and this model is supported

by the evidence in Figure 4. However, in the initial stages

of oxide formation, the weak areas are not necessarily associated

-26-

**!*. T . V- «V:-

' •«»«« - V- ^,-JBl^*.^**- ^ i

'""m

ETCHED REGION

MOIRE FRINGES

PORE INITIATION

iiroiOlnm

Figure 7. A TEM micrograph of the tnitial stages of anodic film formation on 7146 T6. The electron beam is now perpindicular to the interface.

Note that the surface is covered with Moire fringes which may indicate an epitaxial oxide forms during the initial stages of anodization.

-27-

with pores or even regions of oxide thickening. Indeed, the

oxide growth is remote from the dissolution sites. Some pores

contain only a few dissolution sites and some none at all.

Thus, initially the film formation mechanism can be quite

different from latter stages. Further evidence for this is

shown in figure 8, where aluminum dissolution is observed to

occur on the outside of a large pore. In this figure, a

number of well developed pores are also observed to have no

evidence of dissolution at their base; just the opposite, the

base of the pores seems to be filling in with oxide. It is

thought, however, that once the surface is covered uniformly

with oxide, the oxide will redissolve within the pore base.

Very large crystal 1 ographic regions of oxide are apparent in

this micrograph.

At a latter stage of growth, as shown in figure 9, the

pores take a more classical appearance. However, note that

both the inside and outside of the pore walls have a structure

distinct from the interior material. It is thought that the

interior material is removed during the post anodizing treatment.

The mi crostructure of the equilibrium oxide structure

oxide is presented in figure 10. The small pore diameter

present at the initial stage of growth gives way to a pore

diameter characteri Stic of the applied potential and dissolution

now takes place through the base of the pore. This is shown

in this dark field micrograph as a black area in the pore

center. The observations made above are shown schematically

i n f i gure 1 1

.
Ifc*. - ^ If.

I

tr

^ r

1 " 1J
- 28-

DISSOLUTION ON THE OUTSIDE

OF A PORE

I I 5 0 nm

Figure 8. A TEM micrograph showing dissolution of the aluminum around the

outside of an established pre rather than through its center as in the

classica model of Keller, Hunter and Robinson. Note that this behavior

is only observed during the intiial stages of growth.
f
-29-

Figure 9. A TEM micrograph of a latter stage of film formation on 7129 T63

indicating the exi stance of a substructure to the cell walls.

I I 10 0 nm
i ’ll?

f 1

W

* \ 1

•j ;' i;

• ' '•U'
-30-

Figure 10. A TEM micrograph showing the transition from the large numbers

of small pores formed initially to the large diameter equilibrium

pore geometry.

-31-

OXIDE GROWING ON DISLOCATIONS

Figure 11. A schematic diagram

of pore formation on 7000 series

alloys: the initial stage.

Figure 12. A schematic diagram

of a mature pore structure

on 7000 series alloys, (these

models are probably not alloy

dependent)

.
-32-

At a latter stage of growth the structure of the pore

wall seems to change, as does the location of the place in the

substrate where dissolution occurs. A schematic diagram of a

mature pore (cell) structure is shown in figure 12.

3.3. Electrochemical Measurements:

Three representative alloys 7129-T63, 7146-T6, and 99.99

percent aluminum were investigated to determine if observed

differences in their electrochemical behavior could be correlated

with their subsequent adhesion performance.

3.3.1. Current Density versus Voltage Data:

Current density as a function of anodizing voltage curves

are shown in figure 1. From this data, there appears to be

three distinct regimes of growth, marked accordingly in the

figure. Region one, extending from about 0 to 10 volts,

approximates Ohms Law (V = IR) , with a low surface resistance

2 of from 1 to 2 fi/cm . In region two, this resistance increases

2 with increasing voltage to a maximum of around 500 j^/cm ; in

region three, the resistance becomes small. Clearly, in

region three a dominance of the dissolution process over the

film formation process is apparent. In region one, the surface

is relatively free from large amounts of oxide and, in this

unprotected state, is relatively conductive, so its resistance

is low. Film formation and growth takes place in region two.

As the temperature increases, the rate of film formation

increases, along with the rate of film dissolution. In order

to further examine the initial stages of growth, scanning

vol tammagrams of the regions between 0 and 10 volts were made.
-33-

3.3.2. Scanning Voltammetry:

2 Specimens of 1 cm area were prepared for these measurements

in the following way. Each was metal 1 urgical ly polished with

the final step being 0.01 ym MgO. This technique yielded

reproducible measurements which were comparable to those

obtained on el ectropol ished specimens. Simpler grinding

techniques yielded spurious data. Results for 99.99 percent

aluminum, 7146-T6 and 7021-T63 are shown in Figure 13. Each

set of curves corresponds to three successive measurements

starting with a freshly polished surface and zero volts, then

sweeping the potential to 10 volts (limited by our apparatus),

then restarting the sweep. Data from figure 13 is reproduced

i n Tabl e I

.

Tabl e I : Scanning Voltammetry Data.

A1 1 oy 99.99 percent A1 7146-T6 7021 -T63

Break pt. on 1st Sweep

3.1 V/SCE 3.5 2.3

Break on all

successive sweeps

2 V/SCE 21 .9

Resistance before

1st break in curve

2 50 J^/cm 37 54

Resistance

(steady state)

2 18 f^/cm 20 18

During the first three volts, the resistance is markedly

higher than during the higher voltage regions. This observation implies a lack of significant porous film formation

below potentials of three volts/SCE and, probably, pore formation does not start until the first break point, though this
SCANNING

VOLTAMMETRY

CURVES

Zi

PHOSPHORIC

ACID

- 34 -
has not yet been confirmed. Clearly, there is a reproducible

difference in alloy behavior. It would appear that pore

formation in 7021-T63 and in pure aluminum must have similar

forming rates and, probably, similar morphologies, though this

also has not yet been confirmed. This data was obtained in

three percent phosphoric acid electrolytes at a temperature of

60
References

1. Ya. I. A1 exsandrov , Soviet El ectrochemi stry , 1_5, 2(1 979) 141

SEOCAI (2), 137-244 (1979) in English.

2. F. Keller, S. Hunter and D. L. Robinson, J. Electrochem. Soc.

41 1 , (1 953).

3. J. P. O'Sullivan and G. P. Viood, Proc. Royal Soc. London A. 317,

( 1 970)51 1.
-37-

IV. Immersion Deposition;

Many of the results of the characterization of the immersion

deposition process have been published in previous progress reports

and in appropriate journals and will not be discussed again here.

Some results, however, have not been published. These are the

product of scanning voltammetry and simple dissolution experiments

of zinc films in various acids.

4.1. Scanning Voltammetry:

Specimens of 99.99 percent aluminum were metal lographical ly

polished, with the last step being 0.01 ym MgO. The specimen

2 surface area was 1 cm . The voltammetry experiments were

conducted by switching on the potentiostat just after the

specimen was inserted into the electrolyte. The voltage was

swept at a rate of 1 volt/minute to about -2.1 V. Three or

four specimens were run for each set of parameters to insure

that the data presented below is representative

.

In Figure 14a, the curves for the ASTM I electrolyte

(calomel reference electrode) are presented. The Hr^ peak

appeared only on the reverse sweep. The single dip in the

curve was characteri stic of this electrolyte and may be due to

a reduction of the Zn (OH)' ion to zinc. The voltages shown

with arrows are with respect to the hydrogen electrode.

The results of adding ferric chloride and Rochelle salt

to the electrolyte of Figure 14a are shown in Figure 14b. The

original dip in the curve observed in Figure 14a is still

present, however, with at least three more perturbations

apparent. The first, occurring at about 1.44 V(SCE) , may be
H.

SHTC-L2 ZETCATE

ASTM I

i_

-A

- miTS 2.0

j.S^ \ 1 /

^Ih 1 ^

/

< '' /'

/ /

/

/ /

/

/ /

./| 1.5 ,ro

/

0.5

100

! /"

200

MA

Figure 14a. Scanning voltammetry curves for the ASiM I zincate elec

99.99% aluminum.
-39-

/

VOLTS

- 2.0 - 1.8 - 1.6 /I + 1.0

1

ASTM II

SINGLE ZINCATE

99-99^ ALIMINIJM

1 00

ISO

200

250

MA

Figure I4b. Scanning voltammetry curves for ASTM II zincate electrolyte. This

is the same electrolyte as in 1 4a with the addition of ferric chloride and

Rochelle salt.
-40-

I

DOUBLE 2BICATE

99-?9fo ALUI4INUI4

20

VOLTS

1.8 1.6 1.4

//

/

/

20

30

40

50

Figure 14c. Scanning vcl tammetry data for the double zincate procedure with

the same electrolyte as in 14b. This date represents the second zinc immersion.
- 41 -

due to a reduction of Fe( 0 H)

2

~ and implies that the iron would

be deposited before the zinc. This is consistent with results

reported in previous progress reports and by other investigators.

The reactions which occur at voltages greater than 1.8 V (SCE)

may be due to the reductions of iron-zinc, and Rochel 1 e-zi nc

,

compl exes

.

The data obtained for the double zincate procedure are

shown in Figure 14c, and closely resemble the results from the

single zincate procedure from the same electrolyte; i.e.

figure 14b. The main difference observed is in the shape of

the first peak. Instead of occurring at 1.44 V(SCE), as in

figure 14b, it occurs at 1.37 V(SCE). Another peak is clearly

observed at 1.44 V(SCE).

4.2. Dissolution Data:

A large number of experiments were conducted in an attempt

to further improve the zincate procedure. Specimens of about

2 12 cm area were buffed and degreased, etch cleaned, and

desmutted, as in the normal zincate pretreatment procedure.

Most specimens were immersed for one minute in the ASTM III

electrolyte at room temperature unless otherwise indicated.

The specimens were then placed into various acid solutions and

their potential with respect to a calomel or nickel electrode

was monitored as a function of time. A typical voltage-time

curve is shown in figure 15. In most cases, three or four

distinct regions of dissolution are observed. The first part

of the curve, region I, is believed to correspond to dissolution
-42-

Table I; Summary of Dissolution Data.

Regions

A1 1 oy Acid I sec

,

II

. sec

.

III sec

.

Remarks

6061-T6 5?^ H

2 SO4

, 25 °C 14 13 22 Fi gure 2 a , ref. Ni

.

7129-T63 5% H

2 SO

4

, 25 °C 3 15 31 Figure 26, ref. calomel

Zincate at 40 °C.

1100 11 /o H

2 SO4

-.12 16 22 Single Zincate,

Figure 2 c

1100 1 % H

2 SO

4

8 8 23 Double Zincate,

Figure 2 d

1100 1 / 2 % H

2 SO4

17 23 45 Single Zincate,

Figure 2 e

1100 5% H

2 SO4

6 8 12 Figure 2 f

1100 10 % H

2 SO4

2 2 4 Figure 2 g

6061-T6 5% H

2 SO

4

6 11 Fi gure 2 h , ref. Ni

.

7146-T6 ^ 12 -> 21

7146-T6 1 / 2 % H

2 SO

4

8 7 36 Figure 2 i

1/4% H

2 SO4

18 8 >40

7146-T6 3% H

2 SO4

9

20 Figure 2 j

,

3 min. in zincate

5% H

2 SO

4

21

7146-T6 5% H

2 SO4

5 7 26 Double Zincate

9 ir
- 43 -

Time (seconds)

Figure 15. A Typical Dissolution Curve of Zincated

Aluminum
-44-

of the bulk zinc crystallites, and region II to correspond to

a dissolution of the initial zinc layer (see progress reports

III & IV). The rather long plateau, region III, is believed

to correspond to the dissolution of an iron-zinc, or iron

alloy. The data are summarized in Table I.

4.3. Discussion;

The double zincate procedure results in an increase in

the dissolution time of Region III which corresponds to the

dissolution of an iron-zinc alloy. Therefore, the double

zincate procedure is believed to result in enhancement of the

iron alloy coating. This speculation is supported by the

shape of the first peak in the scanning voltammetry data.

Using this dissolution data, specimens were transferred

after Stage I to a sulfamate nickel electrolyte and were found

to form adherent electrodeposits, as was discussed in an

interim report August 1979. This process avoids the use of a

cyanide copper strike but is essentially equivalent to the

ALCOA 661 process.
-45-

V. Adhesion Theory

5.1. Adhesion of Metal ic Coatings*

Whenever two clean metals are placed in intimate contact,

strong adhesive bonding occurs as was shown by Buckley (1). A

theory discussing this type of bond was presented by Ferrante

and Smith (2), Although the theory of Ferrante and Smith was

based on first principles, neither the effect of interfacial

dislocations nor the effect of stored strain energy was included

in their calculations. Several mechanistic theories that have

been published in the literature include the weak boundary

layer theory of Bickerman (3) and a Gri ffi th-Irwin type theory

presented by Good (4) and others (5-9). An understanding of

the forces responsible for bonding are important, not only for

el ectrodeposited coatings, but also for the fields of friction,

wear, and fracture. The development of new coatings for

m.aterials such as aluminum, in which adhesion is usually a

problem, requires an understanding of at least the general

mechanism by which metallic coatings adhere to metals. In

this paper, current theories of interfacial energy are applied

to the calculation of adhesive energy and the specific case of

zinc coated aluminum is discussed.

5.1.2. The Effect of Surface Energy:

As has long been known, the adhesion between two metallic

surfaces is related to the surface energies of each of the two

metals. Consider the total system energy of the metals in

contact and compare it to the total system energy when the
-46-

metals are separated. The energy difference between the two

cases is the adhesive energy per unit area given by.

^ad

“

''c '^s ''i [

1 ]

where and y^ are the surface energies of the coating and

substrate, respectively, and y^. is the interfacial energy.

Equation [1] does not include either plastic deformation or

stored strain energy in the coating or substrate. If two

identical metals of different orientations are considered,

then = y(orientation 1) + y(orientation 2) - y^. [2]

Therefore, in the case of perfect autoepitaxy (y^. = 0), equation

[2] becomes

'ad

= [3]

The relationship given by equation [3] would have to hold

no matter where the interface was located. Therefore, equation

[3], represents the cohesive energy of the material per unit

area. In the case of aluminum, for example, using a bulk

2 value for y

= 840 ergs/cm , the adhesive energy from equation [3]

2 is 1680 ergs/cm . This is only an approximate value because

the surface energy varies with the orientation and measured

values of y are usually an average of many orientations. If two

dissimilar metals are used and y^. is less than Ay where

yc

= y^ + Ay, the adhesive energy would always appear to

exceed the cohesive energy of the material with the lower

surface energy by the amount Ay. Therefore, if a normal force

were applied to the coated substrate, fracture would occur in
-47-

the material with the lower surface energy, and destructive

measurements of adhesion, such as peel tests, would be measurements of the cohesive energies of the lower surface energy

material and would not be a measurement of true adhesion.

In practice, the interfacial energy yi is not <y. Calculations

of adhesive energy must take into account the fact that the

interfacial energy, y^. , depends on the crystal structures of

the material and the coating, the misfit between these two

structures, and the stored elastic strain energy in both the

coating and the substrate. The following calculations for

adhesive energy make use of the theory of interfacial energy

proposed by Jesser and Kuhlman-Wil sdorf (10). Calculations on

the zinc-aluminum interfacial shear modulus given in Appendix

I are based on data in a paper by Rudman and Averbach (11).

In addition, a discussion of the effect of a -pre-existing

crack in the interface on adhesion is presented,

5.1.3. The Effect of Structure:

It is assumed that the potential energy variation experienced by an atom of the coating material moving along the

substrate is given by a Pei erl s-Narbaro function as.

where represents the potential energy amplitude, a^ the

lattice constant of the substrate, and C the potential energy/atom

of a coating atom when sitting in the bottom of the potential

trough. The subscripts x and y refer to the x and y directions.

Under these conditions, the adhesive energy is proportional to

cos + C [4]

a ^
-48-

an average^ potential amplitude, r ’0V = l/2( xoyo V + V ), for N

atoms, Van der Merve has shown (12),

E,d = 2 [5]

where for fee materials K = 3. The potential is related to

G^, the shear modules of the interfaee as was shown by Van der

Merwe as

V 0

G a 0 s

(2Tr)^Nd

[ 6 ]

The subseripts denoting the x and y dependenee have been left

out for elarity, and d is the vertiele separation between the

eoating and the substraet. Using this expression for the

average potential amplitude (6) in equation [5], the adhesive

energy per unit area becomes

^

r\0 s 'ad

27T^d

[7]

Numerical values for the adhesive energies are obtained by

assuming a symmetrical x and y potential distribution and a

close packed configuration with the substrate. The lattice

parameter in the interface is approximated as, proposed by

Jesser and Kuhlman-Wil sdorf (10) by

and the interfacial shear modulus is calculated with the

assumption of a linear correlation of shear modului with

bonding energy (13). Thus, through the atom spacings of the
-49-

coating and substrate, a^ and a^, the effect of orientation

has been introduced into the calculations of adhesive energy.

For epitaxial deposits of zinc on aluminum (14), the adhesive

energies from equation [7] are given by:

A1 (111 )//Zn(0001 )

1530 ergs/cm^

A1 (on )//Zn(0001 ) = 1280 ergs/cm^

A1 (100)//Zn(0001 ) ,E^^ = 930 ergs/cm^

Since the cohesive energy of zi nc is in the neighborhood of

2 1480 ergs/cm , it is seen that failure will occur within the

interface on (Oil) and (100) aluminum planes. This calculation

is only a first approximation and more exact cal ucul ations

have to take into account both the existence of interfacial

dislocations and their effect on bonding, and the elastic

strain energy produced as a result of misfit between the

coating and substrate. In the latter case, the dislocation

density should be reduced. Moreover these calculations represent a minimum adhesive energy. In practice considerable work

would have to be put into the lattice both in the form of

elastic and plastic deformation before planar fracture would

occur.

5.1.4. The Effect of Misfit Dislocations on Bonding Energy:

If a significant misfit between the crystal structure of

the coating and the substrate exixts, a series of interfacial

dislocations will form in such a way as to minimize the resultant

strain energy. The interfacial energy can now be written as
where C, represents, as previously discussed, the potential

energy of an atom sitting in a potential trough. The number

of atoms is given by N and the energy of an array of dislocations

by ( Note : ON = + Y

- 2KMV .)

' S 0

The adhesive energy

can now be written as.

Ead = 2KNV„ -

( E , + E,).

X d yd [10]

The dislocation energy per unit area was shown by Van der

Merwe to be given, leaving out the x and y substrates, as.

1 + 3 - (1 + - 8 In 26 (1 + 8^)^^^ - 26^ [n]

^TT P h

where 6, = — ^ and is a misfit parameter, P is

P(1 - v)(i + Ar

^s^c

the vernier period, a = —, and b is the Burgers vector

« ^c s in the interface. By introducing the effect of dislocations

into the interface, it is seen that the presence of interfacial

dislocation will serve to reduce the adhesive energy over the

dislocation free model. The results of calculations for our

special case of epitaxial coatings of zinc on aluminum are,

A1 (111 )//Zn(0001 ) = 1 300 ergs/cm^

A1 (on )//Zn(0001 ) = 1200 ergs/cm^

A1 (001 )//Zn(0001 ) ;E^^ = 840 ergs/cm^
-51-

5.1.5. The Effect of Strain Energy on Adhesion:

As a result of misfit between the substrate and the

coating, strains will be induced in both the substrate and

coating lattices, the density of misfit dislocations, if they

exist, will be reduced and the adhesive bonding affected in a

complex manner. In the way proposed by Jesser and KuhlmanWilsdorf, the new strained lattice parameters at a strain, e,

are,

a' = a^ + a^e^ and a' = a^ - a^e^ , [12]

c c cc s s ss

and the total energy is given by

E , = 2KNV - E , + E . - E ad 0 X d y d s [13]

It was shown by Jesser and Kuhlman-Hil sdorf (10) that the

strain energy, E^ , can be expressed as

E = G h , s c c

• 1 _ ^

V

c

2 ? o e + e~ + 2v e

- X c

2 2 e + e + 2v e e xs ys xsys

1 - V

[14]

where v is Poisson's ratio and h represents coating thickness

The strain, e^ of the coating is given by

®c

-ba^ ln[2B (1 + - 2 b’^]

G

(1 + (a

2 + a-| )(2iThc)( 1 + v)

[15]
-52-

For a thick substrate, =0. As a result of strain, however,

the dislocation contribution to the interfacial energy will be

reduced. A new value for can be calculated from Equation

[11] by using equation [12] to calculate new values for the

misfit. For zinc coated aluminum the new values of the adhesive

energy are:

A1 (111 )//Zn(0001 )., E^^ = 1300 ergs/cm^

A1 (110)//Zn(0001 ) , E^^ = 1170 ergs/cm^

A1 (001 )//Zn(0001 ) , E^j = 800 ergs/cm^

A comparison between values calculated from the above theory

with results of other investigators is shown in Table I.

5.2. The Effect of a Pre-Existing Crack in the Interface on Adhesion:

A pre-existing crack in the interface can result in

adhesive failure by reducing the stress necessary to propagate

the crack. Such a crack can be a result of poor cleaning,

blistering pores, patches of oxide or other causes. The

Griffith type model considered in this calculation is shown in

figure 16 and consists of a crack of length 2c and unit depth,

located in the interface between a substrate of surface energy

Y-| , and a coating of surface energy y

2

-

The total surface energy of the crack is then given by

Eg = 2 cy^ + 2cy^ = 2c(y^ + Yg) [16]

In order to find the stress, a, at which the crack will

just start to propagate, this energy is equated with the
-53-

Figure 16. A model for a preexisting crack in the interface. The crack is of

unit depth and of length 2C.
-54-

stored strain energy as

E - =2c(y + surface c 's strain

2 2 2 a a ttC E E 2 c s

[ 17 ]

where E represents Young's modulus. Under equilibrium

^ ^ ^surface

~ ^strain^ ^

3c

[18]

so that

In the case of aluminum and zinc, '^critical from equation

[19] is plotted as a function of the crack half length in

figure 17. For crack lengths less than one micron, the stress

9 2 necessary to cause delamination exceeds 2X10 dynes/cm

(30,000 psi). For cracks in aluminum only, the stress (Griffith

Criteria) is 1.95 X 10^ dynes/cm^ (28,300 psi) and for one

micron cracks in zinc 2.2 X 10 dynes/cm (31,900 psi). Note

that once a crack has started, it may run along the interface,

go into the coating or move into the substrate. A more detailed

discussion of crack propagation along an interface was given

by Good (4). In any case, once a crack has started to grow,

the stress necessary to make it continue to grow drops off

very rapidly, so that under conditions of constant stress (and
2500

-55-

NOiivNiiAJvn^a 3snvo

01 AdVSS303N SS3dlS IVOIlldO

Figure

17.

The

critical

stress

required

to

propagate

a

crack

in

the

interface

for

zinc

coated

aluminum
constant crack radius) the coating will delaminate.

5,3. Psuedomorphi sm : the critical radius and the critical thickness*.

The phenomenon of pseudomorphism refers to the tendency

of a coating to assume the crystal structure of the substrate

irrespective of its own crystal structure. The stored strain

energy of a pseudomorphic deposit increases with both the

thickness of the deposit and with the radius of the pseudomorphic

nucleus. Therefore, the deposit is limited in both the radius

and the thickness. The critical radius was given by Cabrera

(1 ) to be

The other quantities in equation [1] were defined in the

previous section. Using the values for zinc epitaxy on aluminum

the critical radius was found to be 4.4 nm or about 20 atom

layers thick. The critical thickness for pseudomorphism given

by Jesser and Kuhlman-Wil sdorf (2) as

R

c

-3b(l + G^a^/4G aj os so

2(1 + GQ/G^)f [

1 ]

where f represents the misfit given by

[ 2 ]

-47t G-j G

2

( 3 2

~ ^1^

2Tra 2 1 ( 1 + z9 ) ( G-j + G^T

He
- 37 -

where for the zinc-a1 umi num system on the (111) aluminum plane

the critical thickness was found to be 0,2 nm or about 1 (at

most) monolayer. 'I'hese calculations seem to preclude this

phenomenon of being of practical importance because crystals

of zinc immersion deposited on aluminum are commonly observed

as being of larger thickness and radius.

Discussion and Conclusions:

Both the existence of interfacial dislocations and the

stored elastic strain energy contribute in a significant way

to the adhesion of metallic coatings on crystalline aluminum

substrates. For non-crystal 1 i ne substrates, aluminum or

otherwise, the adhesive energy for metallic coatings (both

crystalline and amorphous), is higher than for the comparable

crystalline material because of the higher surface energy of

the amorphous material

.

Calculations have been presented only for the case of

zinc coatings on aluminum substrates; however, a basis exists

for future calculations of the relative adhesion between

nickel, iron, copper, tin and various alloy coatings to aluminum,

and results for these materials will be reported in the future.

It was shown above that for zinc and aluminum the greatest

adhesive energy occurs when coatings are deposited on (111)

aluminum compared to the other principle planes. In general

however, the adhesive energy will vary with the surface energy

so that for fee materials the adhesion will decrease in the

following order: ^ ^ ^'lo'^^o^er
-53-

from equation [7], it is predicted that the higher index

planes (greater surface energy planes) would exhibit greater

adhesion than the principle planes. Since extruded parts can

be made in thicker cross section than the comparable rolled

sheet, the incentive of improved adhesion may make a change to

extrusions desirable.

In Table I a comparison between the adhesive energy

determined by a number of different techniques is presented.

The definition of adhesive energy used by Ferrante and Smith

differs from the definition used in this paper by exactly a

factor of two. When allowance is made even for this, the

adhesive energies are still below these calculations by a

factor of about 20 percent. It should be noted that Ferrante

and Smith did not take into account the lattice mismatch

between the zinc and the aluminum, nor did they take into

account the stored strain energy. These factors would have

reduced their values even further.

The surface energy of alloys is in general greater than

the surface energy of pure aluminum and is proportional to the

ultimate tensile strength. It would then seem that adhesion

on alloys would be proportional ly greater than on pure aluminum.

This does not seem to be the case, however. One, therefore,

draws the conclusion that this discrepancy is the result of

either incomplete coverage of the surface, due to incomplete

oxide removal, or defects in the zinc layer resulting in the

existence of a crack. One cause of this, would be imcomplete

coverage over alloying constituents.
Table

1.

Adhesive

Energy

for

Various

Coating -

SubstrateCombinations

(ergs/cm ) - 59 - S IMU

•r“ocJ3

(O cc

<uE

•r—

O)Q.X

LU o oo cT> cn ooIir> oLO

r>.

mm(T) CM

LkJ

00 ^ I— VO n I— I

cn CO f— oCD

cn o to r—

to r—

r— CM CsJ

OroOCO o

CO

-o

tlJi 1 87 85 200

“O CO

CO

00 CO CO

to

to CO CM CM to

UJ CD CD *“ 00

CO I—

a\ f—

LD

CM

CO

CO oC7» CO o CO o VO

CM o CM o r—

CM CM CM CM oO

'JC-M

•t~E

cn-ac

fOOJo

O) to o o to

«3- CO oCM

CO oCM

CO

to

CO

CM

to

LOCM oCM

CO

¥M

>c

OO o o o OO 00 CD CD to LO CO cn to r” 1— r— r— CM

OCX> o(3^ CM

O

OOIjO

Oto

CO

CM o

CM

CM

CT> C

*r“ 4J

«3Oo

a;

4-> <o

&-

4-> OO JO

Z3 cn

r—

C

C h4OO

cr> CO c

CM

CD

O I— c

IMi c

rxi

CD

CD oooo O I— 3o
-6CREFERENCES

1. D. H. Buckley

Proceedings of the Conference on Karlsruhe (1969).

Physics of Adhesion

2. J. Ferrante and J. R. Smith

Sur. Sci. , 2Q, 77 (1973).

3. J. J. Bikerman

J. Coll. Sci. , 2, 163 (1947).

4. R. J. Good

J. of Adhesion, 4, 133 (1972).

5. D. V. Keller

J. of Wear, 6, 353 (1963).

6. M. E. Sikoski

J. of Wear, 7, 144 (1964).

7. H. Krupp

Advan. Col lid Interface Sci., 1_, 111 (1967).

8. R. Eley

Adhesion, Oxford Press (1961).

9. R. E. Baier, E. G. Shafrin, W. A,

Science, 162, 1360 (1964).

. Zesiman

10. W.' A. Jesser and D. Kuhlmann-Wilsdorf

Phys. Stat. Sol., 1£, 95 (1967).

11. P. S. Rudman and B. L. Averbach

Acta Met. , 2, 576 (1954).

12. J. H. Van der Merwe

J. Appl. Phys. , 117 (1963).

13. L. E, Murr

Interfacial Phenomenon in Metals and Alloys

Addison-Wes iey Publishing Co. (1975).

14. D. S. Lashmore

J. ElectrocLem. Soc. , 127 (1980).
- 61 -

LIST OF FIGURES

1. Adhesion as a function of applied voltage for phosphoric acid

anodized 7146-T6 aluminum.

2. The cell and pore diameter as a function of applied voltage.

3. Adhesion as a function of pore diameter.

4. TEM micrograph of 7129-T63 anodized in three percent phosphoric

acid, then plated with nickel.

5. A TEM micrograph of 7129-T63 showing growth of the anodic oxide

along dislocations.

6. TEM micrograph of nickel plated into a porous anodic oxide showing

pore branching.

a. Low magnification bright field image.

b. High magnification bright field image.

c. High magnification dark field image.

7. TEM micrograph of the initial stages of anodic film formation

on 7146-T6 with the electron beam perpendicular to the interface.

8. TEM micrograph showing dissolution of the aluminum around the

outside of an established pore rather than through its center as

in the classical model.

9. A TEM micrograph of 7T29-T63, not nickel plated, and oriented so that the electron beam is perpendicular to the interface. Note

this micrograph is at a latter stage of growth and shows some structure to the cell walls.

10. A TEM micrograph showing a transition between the initial large

number of very small pores formed at low voltage to the larger

diameter pores present at equilibrium.

11. A Schematic Diagram of pore formation on 7000 series alloys.

(Initial stage).

12. A Schematic Diagram of a mature pore structure on 7000 series

alloys (probably not alloy dependent).

13. Scanning voltammetry curves for 60 °C phosphoric acid (3 percent)

on three representative alloys.

14. a. Scanning voltammetry curves for the ASTM I zincate electrolyte

on 99.99 percent aluminum.

b. Scanning voltammetry curves for ASTM II zincate electrolyte

(same as a 14a with ferric chloride and Rochelle salt).

c. Scanning voltammetry data for the double zincate procedure

Csecond zinc immersion).

15 . Dissolution data (a - hj each figure labeled separately.
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N BS-1 1 4A (REV. 9-78J

U.S. DEPT. OF COMM.

BIBLIOGRAPHIC DATA

SHEET

1. PUBLICATION OR REPORT NO.

NBSIR 80-2142

2.Gov't. Accession No. Accession

4. TITLE AND SUBTITLE

Anodizing and Immersion Processes

5. Publication Date November 1980

6. Performing Organization Code

7. AUTHOR(S)

Dr. David S. Lashmore

8. Performing Organ. Report No.

9. PERFORMING ORGANIZATION NAME AND ADDRESS

NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE

WASHINGTON, DC 20234

10. Project/Task/Work Unit No.

11. Contract/Grant No.

12. SPONSORING ORGANIZATION NAME AND COMPLETE ADDRESS rsrreef, City, Sr^re. ZIP)

Aluminum Association, Inc.

818 Connecticut Avenue, N.W.

Washington, DC 20006

13. Type of Report & Period Covered

14. Sponsoring Agency Code

15. SUPPLEMENTARY NOTES

1 Document describes a computer program; SF-185, FIPS Software Summary, is attached.

16. ABSTRACT (A 200'^word or less factual summary of most significant information. If document includes a significant bibliography or

literature survey, mention it here.) This report, summarizes results on the "Plating on Aluminum" project, which was sponsored by the Aluminum Association. The program includes both anodizing and

immersion pretreatments as well as a discussion of adhesion. In addition to the

summary of results, some details of the last seven months of the project are presented

including: (a) the morphology of the initial stages of anodic film formation; (b) the

morphology of the interface of nickel coated anodized aluminum; and (c) a discussion

of adhesion.

17, KEY WORDS (six to twelve entries; alphabetical order; capitalize only the first letter of the first key word unless a proper name;

separated by semicolons)

Adhesion; anodizing; coatings; electrodeposition; immersion deposition; plating;

plating on anodized aluminum

18. AVAILABILITY [^r Unlimitecj

I ] For Official Distribution. Do Not Release to NTIS

I Order From Sup. of Doc., U.S. Government Printing Office, Washington, DC

20402, SD Stock No. SN003-003-

fX~' Order From National Technical Information Service iNTIS), Springfield,

VA. 22161

19. SECURITY CLASS

(THIS REPORT)

UNCLASSIFIED

20. SECURITY CLASS

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UNCLASSIFIED

21. NO. OF

PRINTED PAGES

63

22 . Price

$7.00

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