Effect of Contaminant on the Durability of Epoxy Adhesive Bonds with Alclad 2024 Aluminium Alloy Adherends

SURFACE AND INTERFACE ANALYSIS, VOL. 24, 569-577 (1996)

Effect of Contaminant on the Durability of Epoxy Adhesive Bonds with Alclad 2024 Aluminium Allov Adherends

V

C. L. Olson-Jacques, A. R. Wilson, A. N. Rider and D. R. Arnott DSTO Aeronautical and Maritime Research Laboratory, 506 Lorimer St Fishermen's Bend, Victoria 3027, Australia

Surface preparation of aluminium alloy adherends is designed to minimize the influence of organic contaminant on the durability of epoxy adhesive bonds. The organic contaminant molecules have the potential to inhibit the bonding of the adhesive with the surface oxide on the adherend. X-ray photoelectron spectroscopy (XPS) shows that a freshly exposed aluminium surface quickly adsorbs over ten atomic layers of organic contaminant. Models to fit XPS data measured at various specimen take-off angles indicate that this contaminant is not adsorbed uniformly. Experiments show that for mechanically prepared adherends treated with an organosilane coupling agent, the bond durability depends strongly on the contaminant concentration and the time of administration during the preparation sequence. The observed variations in bond durability can be explained in terms of inhomogeneities in the strength of adhesion between the adhesive and the metal oxide.

INTRODUCTION

The durability of adhesive bonds between 2024 a h - minium alloy plates depends critically on the surface pretreatment applied to the adherend.' A treatment procedure involving a solvent degrease, ScotchbriteB abrade, solvent wipe, grit-blast and an application of coupling agent is known to produce a durable adhesive bond.' The abrasion or grit-blast process affects the adherend roughness and, in principle, removes weak oxide or organic contaminants. However, the introduction of organic contaminant between the abrasion step and the bonding step may adversely affect the strength and durability of the b ~ n d . ~ , ~

Bonded repairs at airbase depots are conducted in an environment with the potential to contaminate the prepared adherend. Typical contaminants include: Avtur fuel aerosols and vapour from operating turbo jet engines; hydraulic fluids and lubricants used in aircraft machinery and water displacing compounds used in corrosion preventative programmes. Volatile solvents used to degrease components may dissolve organic contaminants and then redeposit them on the adherend as the solvent evaporate^.^,^ Whilst polar adhesives can displace and adsorb some organic contamination, and thus make contact with the underlying s~bstrate ,~ the presence of organic contaminant on an adherend can degrade the durability of an adhesive bond.' This paper describes the influence on bond durability of the adsorption of Avtur on the adherend between the grit blast and the bonding step. Avtur (F34) aviation fuel is a mixture of alkanes, including decane C,,H,, (kerosene), de-icing compounds, water insolubility compounds and other additives.

The thickness and distribution of contaminant on the surface is expected to influence the number of successful atomic bonds between the adhesive and the adherend. The surface sensitivity of x-ray photoelectron spectros-

CCC 0142-2421/96/090569-09 0 1996 Department of Defence, Commonwealth of Australia

copy (XPS) provides the ability to measure the thickness and to assess the distribution of thin (< 5 nm) films on a flat solid surface. The change in the attenuation behaviour of photoelectrons as they pass through thin overlayers of contamination can be used in models to calculate the relationship between photoelectron intensity and specimen take-off angle. Although a wide range of models have been published, the limited data available precludes a unique deconvolution without some very restrictive simplifying assumptions. At best, a model can be developed to test a scenario. The laminar model of Bussing and Holloway6 (using a finite number of layers of carbon, each of 0.2 nm thickness, and mean free paths of Tanuma, Powell and Penn') was used to show that approximately ten atom layers of carbonaceous contaminant were present on a polished aluminium alloy surface.* In the present paper, a model is developed using the assumption that the contaminant is distributed across the surface as islands of finite thickness. This model was used to test the assertion that contaminant is adsorbed as discrete islands on the prepared adherends.

Boeing wedge tests (ASTM D3762-79) of bonds between an epoxy adhesive and an aluminium alloy, conducted in warm humid air, indicate that durability loss for bonds with aluminium alloys is generally accompanied by interfacial f a i l~ re . ' , ~*~* ' The rate of bond degradation is determined by the rate of diffusion of moisture to the region just ahead of the crack tip. The changes in bond durability observed with a few monolayers of chemical coupling agents were explained in terms of deformation of the adhesive at the interface with the metal ~ubstrate.~ In this model (Fig. l), applied stress leads to deformation of the polymer at the adhesive metal oxide interface. This deformation will be sen- sitive to inhomogeneities in the strength of bonding and will lead to micro-cavities at this interface. The density of localized strong bonds at the interface will determine the size of the micro-cavities and hence the tortuosity of

Received 14 November 1995 Accepted 22 M a y 1996

570 C. L. OLSSON-JACQUES ET AL.

Figure 1. Model of bond degradation under opening stress showing: (a) large interfacial cavities, leading to rapid moisture diffusion, oxide film degradation and poor bond durability; (b) small interfacial cavities, leading to slow moisture diffusion, slow oxide degradation near the adhesiveloxide interface and good bond durability.

diffusion paths for moisture ingress. The rate of moisture ingress ahead of the crack tip will determine the severity of material degradation at the interface and hence the durability of the adhesive bond. The same model will be invoked to explain durability changes observed for adherends deliberately contaminated with Avtur.

ANGLE-RESOLVED XPS MODEL TO TEST ISLAND DISTRIBUTION OF CONTAMINANT

A simple model consisting of an aluminium substrate partially covered by a carbon contamination layer is considered (Fig. 2). In this model, the photoelectron production is assumed to be constant with depth in the specimen. The detected photoelectron intensity is then calculated by allowing for absorption in the specimen as the photoelectrons propagate toward the detector and performing an integration over the depth of the sample. The detected intensity for C 1s photoelectrons emitted from the islands of carbonaceous material of thickness T , and fractional cover f at an electron take-off angle 8 is

Ic = YPC lSCC Is A,, x (1 - exPC- Tc/&c sin(0)I) (1)

where 5 is the intensity of the x-ray flux, pc is a constant which contains the photoelectron production and detection efficiencies for C Is, Ccls is the carbon concentration and A,, is the absorption half length of the C 1s photoelectron in the carbon overlayer (assuming a

Figure 2. Angle-resolved XPS depicting electrons sourced in the aluminium and the carbonaceous overlayer, showing the attenuation paths for a take-off angle 8, i.e. (90" - 8) to the normal.

simple exponential relationship over all angles for the absorption of the photoelectrons).

The A1 2p photoelectrons are emitted from a semi- infinite solid with some absorption in the fractional carbon overlayer. The intensity for A1 2p photoelectrons is thus

I A i zp(@ = 5 ~ A i zp CAi zp AAA sin(@)

where pA1 2p is a constant which contains the photoelectron production and detection efficiencies for A1 2p, CA1 2p is the aluminium concentration in the aluminium oxide layer, A,, is the absorption half-length of the A1 2p photoelectron in the aluminium oxide and A,, is the absorption half-length of the A1 2p photoelectron in the carbon overlayer. The 1-f term is due to the A1 2p photoelectrons which do not pass through the overlying carbon film, whilst the exponential term allows for absorption of the fraction of A1 2p photoelectrons which pass through the carbon film.

Experimental problems associated with the measure- ment of the I , p and C terms above (which include all production and detection efficiencies within the specimen and equipment) can be avoided by taking the ratios of Eqns (1) and (2) with the 8 = 90" result and then forming the ratio of these results. Thus, the C Is A1 2p photoelectron intensity ratio is

Ic l s ( Q ) / k Is(90") - ( 1 - exPC- Tc/&c sin(8)Il I A ~ z p ( W ~ A i 2,(90") - c 1 - exp( - Tc/&c)l

This equation can be used to fit a curve to experimental data with both Tc and f as fitting parameters. A uniform film result can be obtained by substituting f= 1. Care must be used when using this result for low values off since the results depend on a measurable C 1s signal to form the ratio in Eqn (3).

One limitation of this approach is that it assumes a flat substrate surface and does not allow for any pro- trusion of the carbon film above the surface. At low detector angles the A1 2p photoelectrons that escape from the metal might be masked by either the surface roughness or the carbonaceous material sitting on the surface to give a lower Ai 2p signal than would be pre-

571 DURABILITY OF EPOXY ADHESIVE BONDS WITH ALUMINIUM ALLOY ADHERENDS

dicted by the theory above. This case should be evident through observations of a systematic increase in the C ls/Al 2p intensity ratio at lower detector angles.

The model assumes a simple exponential absorption of the photo-electrons and neglects the effects of elastic scattering and a finite angle of acceptance of the analyser. The latter two effects will introduce a constant error into the C/A1 ratio at take-off angles of up to 20" for a carbon thickness of A/2 and up to 55" for a carbon thickness of A.' The overall effect is a reduction of the substrate counts relative to the overlayer counts, which leads to an increase in T , and thus to an overestimate in the thickness of contaminant.

EXPERIMENTAL

Flat aluminium films formed by evaporation onto glass slides were exposed to Avtur for varying times to produce different concentrations of contamination on the surface. The C 1s and A1 2p spectra from these specimens were recorded at specimen take-off angles of between 90" and 30". These take-off angle measure- ments were conducted with the Kratos X-SAM800 using Mg Ka radiation and operated at a constant pass energy of 20 eV. The theoretical model described in Eqn (3) was evaluated for both uniform and island contaminant distribution.

Adherends for the bond durability tests were cut from 3 mm thick Alclad 2024-T3 aluminium alloy sheet. The standard pre-bond surface treatment used included a degrease with methyl ethyl ketone (MEK), abrasion with a Scotchbrite@ pad, an automated grit-blast with 45 pm alumina gritl0S1l and removal of loose grit with pressurized nitrogen. For convenience, this is defined as the 'Grit-blast' or 'GB treatment. In those experiments where silane coupling agent (SCA) was used, the specimens were dipped in aqueous epoxy silane solution for 10 min and oven dried at 80 "C for 1 h.2 The treatments including the SCA step are designated 'GB + SCA'. The adherends were bonded with Cytec FM73@ epoxy film adhesive using 0.15 mm shims and cured in the press at 120°C for 1 h. Boeing wedge durability tests were conducted in air at 50°C and 98% relative humidity in accordance with ASTM D3762-79.

A series of experiments were conducted in which Avtur (F34) was deliberately introduced during the adherend treatment. In some experiments, designated 'Avtur vap.', the adherends were exposed to Avtur vapour at 50°C and dried. In other experiments, designated 'Avtur dip', the adherends were dipped in Avtur at room temperature for 1 h and then dried. The point of introduction of the contaminant during the adherend preparation process is important and is indicated by a shorthand notation. For example, contamination with an Avtur dip between the 'grit blast' and the application of SCA is written 'GB + Avtur dip + SCA'. A qualit- ative assessment of the bond durability was obtained through Boeing wedge tests described earlier.

The XPS analyses of the fracture surfaces were conducted with a Kratos X-SAM800 spectrometer using Mg Ka radiation, operated in the x 10 lens magnification mode and at a fixed retard ratio of 53. The area analysis was -2 mm diameter, which allowed spectra

to be acquired at points close to the crack front for the wedge specimen. All compositions were calculated using the element sensitivity factors supplied in the Kratos software.

Some 2024-T3 adherends were degreased and then ultramilled without lubricant using a Cambridge Instru- ments diamond blade attachment to a Jung Polycut E microtome. The finish had a sawtooth profile with a 50 pm period and a 2 pm peak-to-valley dimension. These specimens were either bonded with FM 73@ adhesive or dipped in SCA, dried and bonded. Bond durability was assessed with a modified Boeing wedge test using a crack opening displacement of 1 mm. The elastic energy release rate, G, and crack velocity, z), were calculated as described by Venables et ~ 1 . ' ~ Both faces of the separated adherends were examined with a JEOL 6400F field emission scanning electron microscope.

RESULTS

Contaminant distribution by Angle-resolved XPS

The C 1s and A1 2p intensities were measured in the x-ray photoelectron spectrometer at a range of specimen take-off angles for flat films of aluminium evapo- rated on glass slides. The ratio of carbon to aluminium, as a function of take-off angle, is shown in Fig. 3. Expo- sure to Avtur vapour for 15 min and a dip in Avtur for 1 h both increase the carbon film.

Least-squares fits of the data in Fig. 3 to the photoelectron intensity model described in Eqn (3) were conducted with coverage f and contaminant thickness T , as variable parameters. The absorption half-lengths for photoelectrons, A*, and A,,, were estimated at 30 and 26 eV, respectively, based on an approximation to the mean free paths of Tanuma, Powell and Penn.7 More recent estimates of mean free path in carbon for organic compounds13 indicate that the correct selection of mean free path depends on a detailed knowledge of the contaminant and may lead to an increase in the numeric value of the thickness. The values offand T , leading to the best fits, together with the root mean square (RMS) errors, are shown in Table 1. Constraining the fits by setting f= 1, to simulate a uniform contaminant, resulted in contaminant thickness values which were either very low or had no physical meaning (i.e.

Figure 3. Normalized carbon/aluminium ratio as a function of take-off angle 8 measured using XPS.

512 C. L. OLSSON-JACQUES BT AL.

Table 1. Thickness To and surface coverage f of contaminant as a function of exposure to Avtuf

Uniform cover Island cover RMS RMS

Preparation T. (nm) f error T, (nm) f error

Clean -0.4 1.0 0.13 1.5 0.58 0.07 Avtur vap. (5 min) -0.25 1.0 0.09 2.4 0.73 0.03 Avtur dip (60 min) 1.0 1.0 0.38 3.1 0.89 0.05

.The uniform film model gives unphysical film thickness values with Avtur vapour; T, and f were calculated using Eqn (3).

negative). The non-physical nature of these uniform thickness values is compounded by the expected overestimate of the thickness through the use of a model which does not consider the effect of elastic scattering,’ the finite acceptance angle of the analyser,’ or the over- estimation of absorption half-length.I4 Allowing estimates of contaminant thickness and fractional cover to vary in the models, resulted both in positive values of all film thicknesses and lower RMS error values. These results indicate that the contamination is not uniform, with a ‘clean’ surface covered by -60% of adventitious contamination. The thickness and fractional cover of the contaminant both increase with increasing exposure to Avtur. The simple model used here is restricted to islands of a single thickness. In reality, the contaminant is expected to occur in islands of varying size and thickness. However, the extra degrees of freedom introduced into models which include a range of island sizes led to results which were not unique.

Bond durability and fracture surface examination for ultramilled adherends

Figure 4 shows the elastic energy release rate, G, as a function of crack velocity, Y, for Boeing wedge durability tests conducted in condensing humidity at 50 “C for four surface treatments. The value of G generally declines as the crack velocity slows, indicating that the adhesive bonds are degraded by prolonged exposure to the environment under opening loads. The fracture toughness observed for the ‘GB’ adherend preparation is over two orders of magnitude higher than for the ‘ultramill’ preparation. In each case, the organosilane improves the durability of the adhesive bond.

For all samples exposed to condensing humidity at 50°C, failure was adhesive, i.e. at the interface between the adhesive and the metal adherend. Separation pro- duced two surfaces; one surface retained the epoxy adhesive and will be referred to as the side with ‘adhesive’ appearance ; the complementary surface exposed the metal substrate and will be referred to as the side with ‘metallic’ appearance.

The Boeing wedge specimens with the ultramilled adherends were separated to view the fracture surfaces. Electron microscopic examination of the face with metallic appearance shows a meandering crack front extending over - 1 mm (Fig. 5 metal face). A network of very fine aluminium oxide was observed in advance of the crack front and also in the vicinity of dark spots ahead of the crack. These appear to have developed

Figure 4. Elastic energy release rate G vs. crack velocity calculated from Boeing wedge tests conducted in condensing humidity at 50°C for four adherend surface treatments: ‘GB’, ’GB +SCA, ‘ultramill’ and ’ultramill + SCA.

through moisture ingress ahead of the crack front. The generally darker region between the wedge and the crack (subsequently referred to as ‘behind the crack‘) is due to further growth of these oxides in the moist environment present in the open crack, to form hydrated oxide particles. The darkened regions in advance of the crack front (Fig. 5, metal face) corre- sponded with the emergence of voids on the adhesive face (Fig. 5, adhesive face). The dark lines observed crossing the voids are regions where the apex of sawtooth mill profiles of the adherend has penetrated the void and has resulted in an absence of adhesive (Fig. 5, adhesive face). The localized dark regions observed on the metallic face in advance of the crack front are evidence of moisture diffusion to the voids, leading to the formation of hydrated oxide on the metal surface con- nected to the voids (Fig. 5, metal face).

The Boeing wedge specimens in which the ultramilled adherends received SCA treatment were also separated to view the fracture surfaces. Electron microscopic examination indicates an abrupt transition from interfacial failure to adhesive failure in advance of the crack front (Fig. 6). The crack front also meanders over a region of - 1 mm. Voids are clearly seen in the adhesive in advance of the crack front. Behind the crack front, the variation in intensity on the metal face suggests that degradation is not uniform (Fig. 6, metal face). On the adhesive face (Fig. 6, adhesive face), the high magnification image shows a replica in the adhesive of the ultramill profile, with one void breaking through the surface. Energy dispersive x-ray analysis reveals that the light-grey regions near this void contain aluminium, which suggests that moisture from the void has slowly degraded the neighbouring metal. The remaining darker regions on the adhesive face contain only carbon, which is indicative of adhesive. Most of the voids present in the adhesive do not break through to the metal surface,

DURABILITY OF EPOXY ADHESIVE BONDS WITH ALUMINIUM ALLOY ADHERENDS 573

Figure 5. Secondary electron micrographs of the fracture surfaces for the 'ultramill' adherend treatment showing both the metal face and the adhesive face.

which indicates that the SCA assists the adhesive to wet the metal surface during cure.

All subsequent Boeing wedge test data for the 'GB and 'GB + SCA' treatments are shown as plots of crack length against time, since in these cases the 3 mm thick adherends have exceeded the threshold for plastic bending. Therefore, calculations of G and u for these specimens have limited meaning.

Bond durability and fracture surface analysis for grit-blasted adberends

Figure 7(a) shows the Boeing wedge crack growth rates for bonds where the adherends were (i) grit-blasted and (ii) grit-blasted and then treated with the organo-silane coupling agent (SCA). The standard errors for five specimens cut from a single bonded plate are shown as error bars. Results obtained by a number of people over several years are represented as a band. The best case results were carried out at the same time and by the same person as the results shown in subsequent crack growth rate data described in this paper. The application of SCA to grit-blasted adherends clearly decreases

Figure 6. Secondary electron micrographs of the fracture surfaces for the 'ultramill + S C A adherend treatment showing both the metal face and the adhesive face.

the crack growth rate in wedge specimens exposed to condensing humidity at 50 "C, i.e. improves bond durability.

Forced opening of the cracked specimens revealed that failure during exposure to humid air had occurred predominantly along one adhered surface. Analysis of the fracture surfaces in the vicinity of the crack front indicates that for the specimens with the grit-blast ('GB) treatment, the concentration of C, 0 and A1 on the face with adhesive appearance was similar to the concentration of these elements on the face with metal appearance (Fig. 7(b), left). Past experience has shown that the adventitious adsorption of carbon contaminant yields a carbon atomic concentration of N 20%, thus the concentrations observed here do not indicate that adhesive is present on the fracture surface. Hence, this is a clear indication that the fracture has occurred predominantly within the aluminium oxide film on the adherend, as shown diagrammatically in Fig. 7(c) (left), By contrast, analysis of the fracture faces for specimens with the grit-blast and organosilane ('GB + SCA') treatment indicates that the adhesive face is predominantly carbon with very small concentrations of Si, 0 and Al,


Figure 7. Bond durability data and locus of fracture for the ‘GB’ and the ’GB + S C A adherend treatments: (a) crack length vs. root time during exposure to condensing humidity at 50°C showing the band of bond durabilities; (b) fracture surface compositions for both the adhesive face and the metal face measured with XPS; (c) deduced locus of failure.

whereas the metal face has 35 at.% C, plus Si, 0 and A1 (Fig. 7(b), right). The fracture occurs predominantly in the adhesive but is sufficiently close to the SCA layer to show Si and A1 (Fig. 7(c), right). The presence of A1 on the adhesive face may indicate some excursion of the fracture path into the aluminium oxide film.

Figure 8(a) shows the Boeing wedge crack growth rates for bonds in which the adherends were grit-blasted and then exposed to either Avtur vapour or dipped in Avtur. The bond durability with the Avtur vapour exposure was similar to that for the grit-blast (best case) and both were better than that for the Avtur dip.

Analysis of the fracture surfaces in the vicinity of the crack front indicates that for the specimens with the Avtur vapour exposure (‘GB + Avtur vap.’), the composition (Fig. 8(b), left) was similar to that for the grit- blast treatment (Fig. 7(b), left). This is a clear indication that the fracture has occurred predominantly within the aluminium oxide film on the adherend, as shown diagrammatically in Fig. 8(c) (left). For the specimens where the adherends were dipped in Avtur

Figure 8. Bond durability data and locus of fracture for the ’GB’ adherend treatment deliberately contaminated with Avtur: (a) crack length vs. root time during exposure to condensing humidity at 50 “C showing relative bond durabilities for Avtur dip and Avtur vapour; (b) fracture surface compositions for both the adhesive face and the metal face measured with XPS; (c) deduced locus of failure.

(‘GB + Avtur dip’), the analysis also indicates C, 0 and A1 on both fracture faces (Fig. 8(b), right). However, the relative concentration of carbon on the adhesive face had increased whilst that of 0 and A1 had decreased. This suggests that whilst the fracture is predominantly within the oxide film, an excursion of the fracture into organic material may have occurred. One possible rep- resentation is depicted in Fig. 8(c) (right).

Figure 9(a) shows the Boeing wedge crack growth rates for bonds where the adherends were grit-blasted and then exposed to Avtur vapour either before or after the application of organosilane. The bond durability for both cases was similar to that for the ‘GB + SCA’ treatment.

Analysis of the fracture surfaces in the vicinity of the crack front for the specimens with the Avtur vapour exposure between the grit-blast and the organosilane treatment (‘GB + Avtur vap. + SCA’) shows Si on both faces and a very small concentration of A1 on the adhesive face (Fig. 9(b), left). A fracture path which meanders near the SCA layer, as shown in Fig. 9(c) (left), can account for the observed surface compositions. Speci- mens with the Avtur vapour exposure after the


Figure 9. Bond durability data and locus of fracture for the ‘GB + SCA adherend treatment deliberately contaminated with Avtur vapour either before or after the SCA application: (a) crack length vs. root time during exposure to condensing humidity at 50°C showing relative bond durabilities; (b) fracture surface compositions for both the adhesive face and the metal face measured with XPS; (c) deduced locus of failure.

organosilane treatment (‘GB + SCA + Avtur vap.’) had Si on the adhesive face only (Fig. 9(b), right). The concentrations of A1 and C on the adhesive face were slight- ly increased over those for the ‘GB + Avtur vap. + SCA’ treatment. This composition suggests that the fracture may propagate on the metal oxide side of the SCA layer, as shown in Fig. 9(c) (right). The Si composition indicates that there may be less silane present on this specimen.

Figure 10(a) shows the Boeing wedge crack growth rates for bonds where the adherends were grit-blasted and then dipped in Avtur either before or after the application of organosilane. The bond durability for the ‘GB + SCA + Avtur dip’ treatment is similar to that for the ‘GB + SCA’ treatment. By contrast, the ‘GB + Avtur dip + SCA treatment is less durable than the ‘GB’ treatment.

Analysis of the fracture surfaces in the vicinity of the crack tip indicates that dipping the adherends in Avtur between the grit-blast and the organosilane treatment (‘GB + Avtur dip + SCA’) resulted in a very small Si concentration on the metal face only and some A1 on

Figure 10. Bond durability data and locus of fracture for the ‘GB + SCA adherend treatment deliberately contaminated with a dip in Avtur either before or after the SCA application: (a) crack length vs. root time during exposure to condensing humidity at 50°C showing relative bond durabilities; (b) fracture surface compositions for both the adhesive face and the metal face measured with XPS; (c) deduced locus of failure.

the adhesive face (Fig. 10(b), left). This, coupled with the observed carbon concentrations, suggests a fracture which propagates in the organic layer near the oxide surface. The low Si concentration indicates either a very thin SCA layer or a very low surface coverage, as depicted in Fig. 1O(c), left). This interfacial structure is speculative and supporting arguments will be presented later in the discussion section. The composition of the fracture surfaces for the ‘GB + SCA + Avtur dip’ (Fig. 10(b), right) suggests a fracture path which meanders in the organic phase close to the aluminium oxide. One possible structure for the interface and fracture path is shown in Fig. lO(c) (right).

DISCUSSION

In this paper, the focus is on the role of carbonaceous contaminant on adhesive bond strength and durability. The issue that the distribution of carbonaceous contaminant is assessed on flat surfaces, whilst the effect on durability is assessed for the more practical ‘grit-blast’


adherends, needs to be addressed. The elastic energy G released by the adherends to create unit area of fracture surface depends on the roughness of the mechanical surface treatment, the presence of warm humid environment and the presence or absence of a coupling agent (Fig. 4). The SCA improves the durability of an adhesive bond formed with ultramilled adherends, as it does for grit-blasted adherends, but the surface roughness changes G by several factors of ten. It can be reasoned that the adverse effect of carbonaceous Contaminant on a flat surface would be similarly reflected on a grit- blasted surface.

Focusing initially on the ‘ultramill’ treatment, where the adherend surface is relatively smooth (profile with a peak-to-valley of 2 pm and a period of 50 pm), G for the undegraded bond is orders of magnitude greater than the surface free energy (hence work of adhesion) for a1~mina.l~ The majority of the energy is dissipated viscoelastically in processes involving the cooperative actions of large numbers of atoms at the metal oxide/ adhesive interface, thus distributing the stresses at the crack front.16 The rate of degradation of the bond in a moist environment depends on four processes : the hydrolytic stability of the polymer, the hydrolytic stability of the bonds between the polymer and the metal oxide, the hydrolytic stability of the metal oxide itself and the rate of diffusion of moisture to the crack front. For the ‘ultramilled‘ adherend, moisture rapidly degrades the bond (Fig. 4), which implies a rapid rate of diffusion of moisture ahead of the crack front. This rapid diffusion could be facilitated by the formation of interconnecting micro-cavities in the stress field ahead of the crack front3 [Fig. l(a)]. The arguments will be elaborated in the following discussion.

A comparison of the fracture faces of the ‘ultramilled’ and the ‘ultramilled + SCA’ bonded specimens shows that the SCA improves the wetting of the metal by the adhesive. In the ‘ultramilled’ case, the larger voids emerged at the metal surface, leading to degradation of the metal by moisture contained in the voids (Fig. 5). The remainder of the fracture surface in advance of the crack front was generally degraded by moisture ingress. This moisture ingress could be explained by deformation of the adhesive, leading to interconnecting micro- cavities as depicted in Fig. l(a). In the ‘ultramilled + SCA’ case, the voids were inhibited from emerging at the metal surface, through improved wetting of surface by the adhesive. Moisture contained in the few voids which actually emerged at the apexes of the ultramill profile led to some degradation of the metal, resulting in localized areas of aluminium on the adhesive fracture face (Fig. 6). The remainder of the fracture surface (near the crack front) was characterized by the oxide retaining its integrity; i.e. no evidence of aluminium on the adhesive face. Firstly, the SCA inhi- bits the degradation of the metal oxide.3 Secondly, the absence of aluminium on the adhesive face suggests that either polymer hydrolysis or desorption of the SCA and adhesive from the metal oxide surface is the dominant failure process. Thirdly, the observed change in G with crack velocity could be explained by control of the rate of moisture diffusion through a tortuous microcavity structure depicted in Fig. l(b). A change in microcavity structure for the SCA-treated adherends could result from an increase in the density of strong bonding sites

on the adherend due to the improved wetting of the adherend by the SCA. Since wetting of the metal by the adhesive can be influenced by the presence of SCA, the presence of hydrophobic contaminants is expected to adversely influence wetting and hence to assist the diffusion of moisture into the deformed adhesive ahead of the crack front.

Fitting the angle-resolved XPS from flat aluminium surfaces to the model described earlier suggests that the assumption of a uniform contaminant coverage leads to unrealistic film thicknesses (Table 1). Both the ‘clean’ surface and the surface exposed to Avtur vapour appear to be covered with between 60% and 70% of carbonaceous film of >1.5 nm thickness. While island coverage by these contaminants leaves regions of free hydrated aluminium oxide available for bonding to the adhesive, the large fractions of weak surface carbonaceous contaminant are expected to have a more serious effect on the mechanical properties. The pen- etration of the carbonaceous layer by the adhesive during cure, together with the action of the solvents contained in the adhesive, may displace and dissolve contaminant, leading to more effective bonding between the hydrated oxide and the adhesive. Very poor mechanical properties may be expected for the surface dipped in Avtur, since 90% of the surface is covered with a thicker carbonaceous contaminant (Table 1).

The elastic energy release rate G depends on the roughness of the mechanical surface treatment (Fig. 4). Some of the increase for that of the ‘GB over the ‘ultramill’ treatment may be attributed to an increase in surface area: some may be attributed to an increase in the efficiency of surface wetting by roughening the surface;” and some may be attributed to the influence of local stress fields in the vicinity of the micro- roughened grit-blast surface. Each of these factors will have an effect on the structure of micro-cavities formed ahead of the crack front and hence on the rate of diffusion of moisture into this region. The presence of SCA or of hydrophobic contaminant will have an effect on wetting of the surface and may have an effect of the distribution of stress in the polymer in the vicinity of the grit-blasted adherend.

Exposure of grit-blasted adherends to Avtur vapour above a liquid fuel at 50 “C has minimal effect on both the initial crack length in dry conditions and bond durability [Fig. 8(a)]. For both the ‘GB’ and the ‘GB + Avtur vap.’ adherend treatments, the fracture path remains within the oxide film, indicating that the adhesive has effectively counteracted the presence of surface contamination by the Avtur vapour. For the ‘GB + Avtur vap. + SCA’ and the ‘GB + SCA + Avtur vap.’ treatments, again the introduction of Avtur vapour has minimal effect on both the initial crack length in dry conditions and the bond durability [Fig. 9(a)]. The fracture path shows some minor shifts toward the oxide film with the introduction of Avtur vapour (of Fig. 9(c) and Fig. 7(c), right). This shift is indicative of very minor weakening of the SCA interface by the contaminant, but the effect is not sufficient to be reflected in a change in bond durability.

Dipping grit-blasted adherends in liquid Avtur fuel at room temperature has a small effect on bond durability [Fig. X(a)]. However, the weakening of the interface is not commensurate with the expected 90% coverage


with contaminant (Table 1). The adhesive appears to dissolve some contaminant during cure, leaving a small residual as depicted in Fig. 8(c) (right). The presence of the carbonaceous contaminant will locally weaken the adhesive/metal oxide interface and the fracture path makes an excursion from the oxide into carbonaceous material. Under load, a small change in the extent of the submicroscopic cavities (Fig. 1) will lead to an increase in moisture diffusion ahead of the crack front and thus to the small decrease in bond durability. The photogra- phs shown in Fig. 5 indicate that whilst voids in the adhesive contribute to moisture diffusion and localized attack of the metal substrate, diffusion at the bond interface leading to generalized metal oxide interface degradation ahead of the crack tip is important. This degradation is supported by evidence of the growth of weak oxide particles in the high-magnification image shown in Fig. 5. The meandering form of the crack front provides evidence of inhomogeneity in the strength of bonding between the metal oxide and the adhesive under load, which is a feature required by the model shown in Fig. 1.

For the ‘GB + Avtur dip + SCA’ treatment, the dip in Avtur before the application of SCA leads to a mar- ginally lower bond durability than the ‘GB treatment and a substantially lower durability to that of the ‘GB + SCA’ treatment [Fig. lqa)]. The aqueous epoxy silane solution appears to be ineffective in displacing the contaminant i.e. is ineffective in wetting the contaminant surface. The contaminant is exposed to the drying cycle at 80°C for 1 h before the adhesive is applied. Whilst the adhesive may dissolve some of the organic material during cure, some residual hydrocarbon fractions are expected to remain. The characteristics of the failure surface composition [Fig. 10(b), left] can thus be explained in terms of a mixture of adhesive, SCA and contaminant influencing the interfacial bonding [Fig. lO(c), left]. The durability of the ‘GB + Avtur dip + SCA’ [Fig. lqa)] appears to be controlled by the cohesive stability of the adhesive [Fig. lO(c), left], whereas the durability of the ‘GB’ treatment [Fig. 7(a)] appears to be more dependent on the hydrolytic stability of the aluminium oxide [Fig. 7(c), left].

For the ‘GB + SCA + Avtur dip’ treatment, the dip in Avtur after the application of SCA leads to a minimal

decrease in bond durability [Fig. lqa)]. Again, the solvents in the adhesive are expected to be effective in displacing and dissolving the contaminant. The fracture path [Fig. lqc), right] indicates that the SCA retains substantial strength, although there is evidence of the onset of moisture ingress leading to some degradation of the oxide.

CONCLUSION

Avtur is adsorbed on a flat aluminium surface as islands of at least 70% coverage for vapour exposure and at least 90% for a dip in liquid fuel. The presence of this contaminant does not necessarily degrade bond durability. Application of the Avtur just before the adhesive is applied has little effect on bond durability. However, dipping the adherend in Avtur before the application of the epoxy silane coupling agent leads to a dramatic reduction in bond durability.

Fracture of bonds in condensing humidity at 50 “C is adhesive, i.e. between the adhesive and the metal adherend. The grit-blast treatment leads to a fracture path predominantly in the oxide film on the adherend. Treat- ment with the epoxy silane coupling agent shifts the fracture toward the polymer. The introduction of contaminant can lead to subtle shifts in the fracture path which reflect local changes in the relative hydrolytic stability of the polymer, the adhesive to adherend bond and the metal oxide. The shifts in fracture path are sig- nificant for the contaminanation adsorbed on the adherend during a dip in Avtur fuel. The fracture behaviour in these cases suggests that bond durability is controlled by the relative dominance of polymer hydrolysis, polymer desorption or oxide degradation. The rate of moisture ingress into the stressed adhesive ahead of the crack front can account for the observed changes in bond durability.

Acknowledgements

Mr John van den Berg is thanked for his work in preparing the Boeing Wedge Tests.

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