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Mechanical and frictional behaviour of nano-porous anodised aluminium
Materials Chemistry and Physics 148 (2014) 887e895
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Mechanical and frictional behaviour of nano-porous anodised aluminium N. Tsyntsaru a, *, B. Kavas b, c, J. Sort d, M. Urgen b, J.-P. Celis e a
Institute of Applied Physics of ASM, 5 Academy str., Chisinau, MD 2028, Moldova Istanbul Technical University, Department of Metallurgical and Materials Engineering, 34469 Maslak, Turkey c Ford Otomotiv San A.S., Istanbul, Turkey d Catalana de Recerca i Estudis Avançats (ICREA) and Departament de Física, Universitat Auto noma de Barcelona, E-08193 Bellaterra, Spain Institucio e KU Leuven, Dept. MTM, Kasteelpark Arenberg 44, B-3001, Belgium b
h i g h l i g h t s
g r a p h i c a l a b s t r a c t
Well-ordered porous AAO with pore diameters between 16 and 75 nm were produced. Porosity and composition of electrolytic baths influence the mechanical properties. Ball-on-flat configuration was used in tribological testing under dry conditions. Adherent tribolayer consisting of fine worn particles prevents AAO from cracking. Testing parameters are moreover essential to envisage AAO practical applications.
a r t i c l e i n f o
a b s t r a c t
Article history: Received 13 March 2014 Received in revised form 28 August 2014 Accepted 30 August 2014 Available online 17 September 2014
The porous structure of anodic aluminium oxide (AAO) can be used in versatile applications such as a lubricant reservoir in self-lubricating systems. Such systems are subjected to biaxial loading, which can induce crack formation and propagation, ultimately leading to catastrophic mechanical failure. In this study, the mechanical and tribological behaviour of AAO, prepared from two different types of electrolytes (sulphuric and oxalic acids), are studied in detail. The electrolytic conditions are adjusted to render highly tuneable average pore diameters (between 16 and 75 nm), with porosity levels ranging from 9% to 65%. Well-ordered porous AAO are produced by two-step anodization at rather low temperatures. Mechanical properties, mainly hardness and Young's modulus, are investigated using nanoindentation. Both the porosity degree and the composition of the electrolytic baths used to prepare the AAO have an influence on the mechanical properties. Ball-on-flat configuration was used to estimate the tribological behaviour under dry conditions. No major cracks were observed by scanning electron microscopy, neither after indentation or fretting tests. In the running-in period of tribology experiments the pores
Keywords: Nanostructures Oxides Coatings Mechanical properties Tribology and wear
* Corresponding author. E-mail addresses: [email protected], [email protected] (N. Tsyntsaru), [email protected] (B. Kavas), [email protected] (J. Sort), [email protected] (M. Urgen), [email protected] (J.-P. Celis). http://dx.doi.org/10.1016/j.matchemphys.2014.08.066 0254-0584/© 2014 Elsevier B.V. All rights reserved.
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were filled with debris. This was followed by the formation of a highly adherent tribolayer e a third body consisting of fine worn particles originated from both the sample and the counterbody. Pore diameter and porosity percentage are found to strongly affect hardness and Young's modulus, but they do not have a major effect on the frictional behaviour. © 2014 Elsevier B.V. All rights reserved.
1. Introduction The unprecedented growth in different areas of nanotechnology has prompted the rapid development of novel routes (advanced lithography procedures, wet chemistry, self-assembly, etc.) to prepare materials with reduced lateral dimensions. The characterization of anodic aluminium oxide (AAO) is of outmost importance since this material has possible applications in fields such as catalysis, chemical sensors, biosensors, filters, humidity sensors, as well as ceramic materials for implant materials [1e5]. The AAO can also serve as templates for the growth of electrodeposited nanowires [6] and as a support for measuring the mechanical properties of nanocarbon tube ropes [7], amongst others. AAO is also potentially interesting for tribological applications since the nanoporous structure can be used as a reservoir or template for solid lubricants (e.g., nano-tubes or nano-fibres) to form self-lubricating structures [8]. AAO nanostructures are typically composed of two regions: a relatively pure alumina inner layer, consisting entirely of Al2O3, and an acid anion-contaminated outer layer, resulting from the incorporation of anions into the alumina structure during anodizing [9]. The presence of these two regions together with the occurrence of different crystallographic phases (e.g., amorphous, a-Al2O3 and gAl2O3, that stem from different heat treatment conditions) affect the mechanical strength. Friction and wear performance of filled-in AAO have been studied to some extent [10]. Fundamentally, homogenous and heterogeneous (i.e., composite) ceramics behave in different ways when subjected to normal loading. Namely, homogeneous ceramics typically respond to the intense confined shear under the sharp indenter tip by an elastic deformation followed by the formation of radial cracks due to elastic mismatch. Conversely, heterogeneous ceramics can accommodate elastic strains much more efficiently, hence minimizing crack propagation and showing improved fracture toughness [11]. In spite of the previous studies on the mechanical properties of AAO [12e16], there is still lack of knowledge in some particular issues, such as the detailed understanding on how the porous structure of AAO and the concomitant mechanical properties (hardness, Young's modulus) evolve during mechanical deformation. Information concerning the tribological properties of AAO templates is also rather limited. Some literature exists on the tribological characterization and friction behaviour of AAO filled up with solid or liquid lubricants [17e19]. However, the frictional behaviour of AAO itself has been poorly investigated yet. In this study, we investigate (i) the mechanical properties of AAO using nanoindentation with a Berkovich indenter tip, and (ii)
the tribological response in a ball-on-flat sliding tester. The effect of porosity level, pore diameter, and the type of anodizing electrolyte on mechanical and tribological properties are described. Particular emphasis is given on how the mechanical tests modify porosity and, consequently, on how they affect the mechanical properties of AAO. The acquired understanding of the evolution of the physical response and the wear behaviour of AAO under mechanical loading is of crucial importance in many mechanical systems, and triggers new ideas for innovative application areas. 2. Experimental 2.1. Surface preparations and anodizing Al sheets with a thickness of 0.25 mm (99.99%, Alfa Aesar Johnson Matthey GmbH) were electropolished at a constant current density of 500 mA cm1 for 1 min in an electrolyte consisting of HClO4 (60 wt. %) and C2H5OH (abs.) in a volumetric ratio of 1:4. Temperature was kept constant at ~10 C. Samples were examined using scanning-electron microscopy (SEM) and non-contact white light interferometer (WYKO) before and after electropolishing for the estimation of Ra (arithmetical mean roughness) and Rz (tenpoint mean roughness) values. Surface roughness decreased after electropolishing (i.e. Ra from 740 nm to 250 nm, and Rz from ~7 mm to ~3.2 mm). The anodizing of electropolished samples was carried out in two different electrolytes, namely 20 wt.% H2SO4 [20] and 0.3 M H2C2O4 [21], in order to produce AAO with different characteristics (e.g., pore size or interpore distance). The synthesis procedures used are described in Table 1. The samples were anodised for various electrolysis durations, so as to obtain AAO with an average thickness of ~17 mm. Two-step anodisation was successful, as can be seen from the SEM images (Fig. 1). Well-ordered, hexagonally distributed pore arrays were obtained with different pore diameters and porosity degree (Table 2). In order to simplify the nomenclature of the samples, samples produced in sulphuric acid are denoted hereafter as S1 and S2, respectively, while those anodized in oxalic acid and pore widened afterwards are identified as O1, O2 and O3 (Tables 1 and 2). Pore widening was performed at room temperature (~25 C) by immersion of anodized in oxalic acid AAO in H3PO4 (5 wt.%) for different durations namely 32 min (O2) or 75 min (O3) to obtain variable pore diameters. Similar pore diameters were produced for samples O1 and S2 in order to assess the influence of the type of electrolytic bath on the wall thickness and the mechanical
Table 1 Overview of the samples preparation conditions. Sample name
Anodizing, I step
S1 S2 O1 O2 O3
15 V Sulphuric 500 rpm, 1 C, 10 min 21 V Oxalic 40 V, 800 rpm, 25 C, 60 min
Oxide removal 10 min 15 min 60 min
CrO3 (1.8 wt. %) þ H3PO4 (6 wt. %) 60 C
Anodizing, II step
Pore widening
15 V, 97 min Sulphuric 500 rpm, 1 C 21 V, 20 min Oxalic 70 V, 800 rpm, 5 C, 60 min
No No 32 min 75 min
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properties. Chemical composition of AAO was quantified by energydispersive spectroscopy (EDS). The EDS analyses revealed the presence of carbon and sulphur contaminations in the AAO samples produced from oxalic and sulphuric acids baths, respectively. 2.2. Mechanical and tribological characterization To study the mechanical behaviour of anodized aluminium oxides, nanoindentation tests were performed at normal loads varying from 5 to 100 mN using an indenter from CSM Instruments, equipped with a Berkovich-type tip. Loading rates were chosen to be twice the value of the peak applied load per minute. Indentation imprints were observed by SEM. Loading vs. penetration depth data were recorded. Hardness values were evaluated using the method of Oliver & Pharr [22] from a total of 5 indentation curves. Wear behaviour was investigated in relation to pore diameter, porosity percentage and oxide composition. Tribological tests were performed in the KUL-MTM fretting mode I apparatus [23] and a Modular Universal Surface Tester (MUST, Falex Tribology N.V., Belgium) in a ball-on-flat configuration operated under reciprocating mode. The counterbody was a 5 mm diameter corundum ball reciprocating on AAO samples at a frequency of 1 Hz, a displacement of 100 mm, under normal loads of 40 mN, 80 mN (mesoscale) and 1000 mN (macroscale). The mesoscale and the macroscale terms are derived from the displacement and normal loads used as is given e.g. in Ref. [24]. The test temperature and humidity were kept constant at 23 C and 50% RH, respectively. It should be noted that steel balls had been used as counterbody in previous works [25], and this resulted in tribochemical reactions during sliding. Here, to avoid these effects, corundum balls were used as counterbody, i.e., the same material as the base membranes (AAO). Samples were degreased before the reciprocating sliding tests. After sliding tests, the samples were ultrasonically cleaned in ethanol for 7e10 min to remove the debris. 3. Results and discussion 3.1. Mechanical characterization of nanoporous AAO by nanoindentation
Fig. 1. Scanning electron microscopy (SEM) images of produced anodised aluminium oxides (AAO).
SEM images of the nanoindentation imprints are shown in Fig. 2. Since AAO is a ceramic (i.e., non-ductile material), crack formation and propagation could be a priori expected as a result of the indentation compressive and shear stresses. However, SEM images revealed that the investigated AAO did not show major cracks either inside or near the imprint, as opposed to some previous studies [14]. The lack of major cracks might be related to the intrinsic structure of the investigated AAO or to the difference in strain rate and other test parameters, as compared to the studies carried out by other authors. Similar results with our tests were obtain in [26], but with less evidence of different behaviour between samples anodised in sulphuric and oxalic acids. The anodized layers show differences in structure when compared to bulk brittle Al2O3. They contain substantial amounts of OH groups, as was previously evidenced [27]. These OH groups cause an increase of the overall intrinsic plasticity, thus inhibiting crack nucleation. The high-magnification images (see Fig. 2, insets) reveal that minor cracks are actually formed within the imprinted area during the course of the indentation experiments in samples S1 and S2. Sulphur content and the structure of AAO produced in sulphuric acid might have made the oxide more brittle. On the contrary, samples O1, O2 and O3 do not show any sign of cracking. For given indentation conditions (i.e., maximum load) the number of cracks formed as a result of indentation was less on sample S2, which contained wider pores than sample S1. This suggests that the
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Table 2 Overview of the morphological features of the investigated samples. Sample
Pore diameter, nm
Interpore distance, nm
Porosity, %
S1 S2 O1 O2 O3
~16 ~27 ~27 ~45 ~75
~39 ~56 ~94 ~94 ~94
~15 ~20 ~9 ~20 ~65
pore diameter affects the formation of cracks. When comparing O1 and S2, which had both approximately the same pore diameters, we did not observe any cracks on O1. This may be related to the differences between the structural composition of AAO films produced in sulphuric and oxalic acids. No cracks were observed either on samples O2 or O3. Loadepenetration depth curves were measured on all samples. As an example, Fig. 3 shows representative indentation curves obtained on sample O3 for maximum applied loads of 5, 20, 50 and 100 mN. As it was noticed on all loadingeunloading curves, no popin events corresponding to major cracks or phase transformations took place during the nanoindentation experiments. The indentation curves were rather smooth without any obvious sudden changes. The average values of maximum and final penetration depths, as well as the elastic recovery percentages, were estimated from the indentation curves (see Table 3). The results indicate that the elastic recovery of AAO produced in sulphuric acid, regardless of their pore diameter or porosity degree, is larger than for samples produced in oxalic acid. Indeed, although O2 and S2 show similar porosity percentages (Table 2), a significant difference in the elastic recovery is observed (i.e., while O2 has a recovery of ~24%, that of S2 is ~41%). Similarly, although the pore diameters of samples O1 and S2 are similar, a clear difference is again observed on the elastic recovery. Remarkably, samples produced in oxalic acid electrolyte show similar recovery percentages in spite of their different pore diameter and porosity level. This seems to suggest that, rather than the porosity degree and pore size, the type of electrolytic bath used for anodizing plays a major role in the elastic recovery of the samples. Contrary to what is observed for the elastic recovery, the porosity level is found to have a strong influence on the hardness and reduced Young's modulus. Namely, the maximum penetration depth attained in sample O3 (with a porosity of ~65%) is much larger than on samples O2 (with porosity of ~20%) or O1 (with porosity of ~9%). As a consequence, sample O3 is mechanically softer (see Fig. 4). The effect of porosity on the Young's modulus of the AAO porous films was also analysed. As shown in Fig. 5, there is a clear trend for Er to decrease with the increase of porosity level. This is particularly clear for the oxalic samples, especially at low indentation loads, where Er decreases from 110 GPa (for 9% porosity) to 103 GPa (for 20% porosity) and to 60 GPa (for 65% porosity). The reduced Young's modulus takes into account the elastic displacements that occur in both the specimen, with Young's modulus E and Poisson's ratio n, and the diamond indenter, with elastic constants Ei ¼ 1140 GPa and ni ¼ 0.07. The relationship between Er and E can be expressed as [28]:
1 1 v2 1 v2i þ ¼ Er E Ei
(1)
For bulk Al2O3, the Poisson's ratio is around 0.24 [29]. This value could be slightly lower if one takes into consideration the existence of porosity [30]. Assuming that n z 0.24, the Young's modulus of the porous Al2O3 would be: E z 114 GPa for 9% porosity, 106.6 GPa
Fig. 2. Scanning electron microscopy (SEM) images of the AAO templates indented using a maximum load of 100 mN. Shown in the insets are high magnification images of the indented regions.
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Fig. 3. Loadingeunloading nanoindentation curves corresponding to sample O3 for maximum applied loads of 5, 20, 50 and 100 mN.
for 20% porosity and 59.7 GPa for 65% porosity, hence in practice very similar to Er. The porosity level is known to have a strong influence on the elastic constants of metallic and ceramic materials. In a first approximation, it has been shown that [31e33]:
Eporous ¼ Ebulk
rporous n rbulk
(2)
where rporous/rbulk is the relative density. The exponent n has been reported to be equal to 2 for open-celled foams or sponge-like structures [31,32], whereas n ¼ 3 is generally assumed for materials whose porous structure is arranged forming a honeycomb hexagonal array normal to the surface [33]. In the case of porous Al2O3, values close to n ¼ 1 have been reported [33]. The relative density is related to the porosity volume fraction:
rporous ¼1P rbulk
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Fig. 4. Dependence of hardness on the maximum penetration depth for the investigated AAO samples.
Concerning the hardness tests (Fig. 4), the values obtained for samples O1, O2 and O3 are also lower than those reported for nonporous anodized alumina, which have been reported to be around 13 GPa [34]. The decrease of hardness (or compressive yield stress) with porosity is also a well-documented effect [35e37] and has been modelled using finite element simulations of nanoindentation curves [36]. An equation analogous to Eq. (2) may be used to correlate the yield stress of the porous structure with that of the bulk solid material:
sporous ¼ C2 sbulk
rporous rbulk
m (4)
where C2 ¼ 0.3 and m is an exponent [37]. Although the relationship between hardness and yield stress in bulk oxide ceramic is
(3)
The Young's modulus of bulk (non-porous) amorphous anodized Al2O3 (non-heat treated) has been reported to be around 147 GPa [34]. Using Equation (2), values of n equal to 2.8, 1.5 and 0.85 are obtained for indentations made using a maximum load of 5 mN, for samples O1, O2 and O3 with 9%, 20% and 65% porosity, respectively. These values of n fall within the range of expected values for porous materials. The values of n remain similar for samples O1 and O2, regardless of the applied indentation load. Conversely, for sample O3, the exponent n increases with the applied load, up to n ¼ 1.2, probably due to the densification effects observed during nanoindentation (see SEM image in Fig. 2).
Table 3 Summary of the values of maximum depth, final depth and elastic recovery for the different AAO templates, after indentation tests using a maximum load of 100 mN. Sample
Max. penetration depth, nm
Final depth, nm
Recovery %
S1 S2 O1 O2 O3
~1320 ~1325 ~1010 ~1110 ~1650
~840 ~780 ~710 ~845 ~1260
~36 ~41 ~29 ~24 ~24
Fig. 5. Dependence of the reduced Young's modulus on the maximum penetration depth for the different investigated AAO membranes.
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often taken as Hbulk z 1.6sbulk [38], a quantitative relationship between hardness and yield stress in porous materials has not been readily established yet. Assuming that Hporous ¼ 1.6sporous [38], Eq. (4) would give m z 0.4 for the samples with 65% porosity and applied load of 5 mN (and lower values for the samples with lower porosity degree), which are much smaller than the exponent obtained in an open-cell metallic foam (m ¼ 1.5) [37], but in agreement with the decrease of hardness with porosity observed by other authors on anodized alumina [34]. The samples S1 and S2 show relatively similar hardness values at all applied loads, varying between ~4.5 and ~6 GPa, except at small loads where there is large scatter (see Fig. 4). The Young's modulus is also similar for these two samples (Fig. 5). This is probably due to the very small difference in the porosity percentage between these two samples. The unexpected larger hardness recorded on sample S1 at 5 mN could be linked to the rougher surface of this sample. 3.2. Frictional behaviour of nanoporous AAO at meso- and macroscale loading Nanoindentation tests provide evidence of the range of normal forces that AAO templates can withstand without being catastrophically damaged. In analogy, the evaluation of the tribological behaviour is important to assess the range of forces that can be applied under sliding to this material before it undergoes dramatic frictional effects. In order to evaluate how the porous structure behaves when subjected to bidirectional loading, two types of reciprocating sliding cycles were set: (i) 50 cycles to reveal the initial evolution of wear mechanism and (ii) 500 cycles to observe the frictional behaviour already in the steady-state regime. Wear tracks were examined using SEM after 50 cycles performed at relatively low loads (40 and 80 mN) (Fig. 6). The observations revealed an interesting response; namely, pores started to be filled-up with fine debris particles without any cracking. This behaviour proved that the porous tubes were strong enough to resist bidirectional loading at this level of applied forces. Tangential force versus displacement hysteresis loops showed that the tests were carried out under gross slip regime. In the running-in period of meso scale tests (i.e., at applied load of 40 and 80 mN), the tangential force was first rather low and increased with the number of cycles (Fig. 7). This demonstrates that abrasive wear at the beginning causes the formation of wear particles which act as third body making the relative motion more difficult due to adhesion. As a result, the wear mechanism is a combination of abrasive and adhesive wear, hence resulting in a high coefficient of friction (Fig. 7). At the steady state, the coefficient of friction reaches approximately 0.9. For higher
Fig. 7. Evolution of the coefficient of friction for S2 with the number of cycles under different applied loads.
loads (i.e., macro-scale loading at 1000 mN), the running-in period is followed with a decrease of the coefficient of friction until a value close to 1 is reached in the steady-state (Fig. 7), quite close to that measured under meso-scale loading. Increasing normal load from 40 mN (Fig. 8a) to 80 mN (Fig. 8b) under the same displacement amplitude brings about a decrease in effective sliding distance in the contact with sliding cycles. This can also be linked to the adhesion of debris or the penetration of the ball into the wear track. As at macro scale loading (Fig. 8c), tangential force shows a decrease with the number of reciprocating sliding cycles. The difference between the evolution of hysteresis loops under 1000 mN and meso-scale loading can be linked to the penetration depth of the counterbody which was deeper at higher normal loads. Once the ball creates a wear track by abrasive wear, the relative motion becomes easier, revealed by a decreasing tangential force. The stiffness of the macro-scale test machine is larger than that of meso-scale tester [26]. Thus, the obvious effect of adhesive wear observed at meso-scale loads is less apparent at macro-scale loads. Evolution of the coefficient of friction with the number of cycles for all samples, at 1000 mN normal load, is given in Fig. 9. A high coefficient of friction was expected for all samples since a corundum counterbody was used (i.e., the same material as the test
Fig. 6. Scanning electron microscopy (SEM) images of S2 samples after 50 cycles fretting under 40 mN (a) and 80 mN (b).
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Fig. 8. Fretting hysteresis loops registered for S2 under 40 mN (a), 80 mN (b) and 1000 mN.
sample). Similar coefficient of friction dependencies were noticed on samples O2, S1 and S2 (which had rather similar porosity percentages). The coefficient of friction recorded on these samples was ~1.35 at the beginning and decreased down to ~1.1 once steady state sliding is established. The resemblance of the results can indeed be expected from the similar porosity percentages, which allow analogous amounts of debris entrapped in the pores and also result in similar penetration depths. The coefficient of friction recorded on O1 is higher than on S1, S2 and O2 probably due to a lower porosity which means bigger amount of aluminium oxide in the contact area. In this sense, sample O3, with the highest porosity level, should exhibit the lowest coefficient of friction, which was not noticed during our experiments. The reason for the relatively high coefficient of friction (~1.5) in this highly porous sample could be the high penetration depth achieved on O3 (similar to what was observed during nanoindentation tests). As the ball goes deeper inside the oxide layer, it becomes surrounded by a larger amount of nanotubes as if it became buried within the AAO. As a result, higher tangential force is required to carry out the reciprocating motion. Furthermore, this sample is also able to entrap larger amounts of debris particles and retards the formation of a top-layer consisting of debris that induces a decrease in the coefficient of friction. SEM images of wear tracks produced on S2 sample after 500 cycles are shown in the upper row of Fig. 10. Debris was sparsely
Fig. 9. Evolution of the coefficient of friction of all samples under an applied load of 1000 mN.
distributed on the surface in piles forming a third-body tribolayer. The tribolayer consisting of compacted wear particles exhibits an encrusted structure. The debris adhering locally at some parts in the wear track causes a complex wear mechanism stemming from a combination of abrasive and adhesive wear together with a high coefficient of friction (Fig. 9). This debris is highly adherent and remains visible even after ultrasonically cleaning the samples for 10 min in ethanol. Wear tracks on sample O2 were similar to those produced on S2 (Fig. 10). Compacted debris particles inside the wear track and a layered structure with sparsely distributed debris were observed. Indeed, analogous complex wear mechanism as that on S2 is expected on O2, thus resulting in a high coefficient of friction. On sample O2 pores partially filled with debris in areas where compacted debris did not adhere, were still visible after 500 cycles. Conversely, the wear tracks on the highly porous sample O3 show some different features (Fig. 10). The light greyish areas observed on samples O2 and S2 at low magnification were not present on sample O3. Moreover, at higher magnification, the layered and compacted debris were neither as dense nor as thick as those seen on samples S2 and O2. Pores remained more visible. Because of the larger pore size, the nanotubes were easily filled with most of the debris. Such progressive filling probably inhibits, to some extent, the formation of a dense and thick top layer. Nevertheless, as can be envisaged from the wear tracks, the nanotubes were just overfilled after 500 cycles and in case of running the reciprocating sliding tests for longer durations, similar top layers as on samples S2 and O2 would probably form. Although the tubular structure of sample O3 was not significantly damaged, crack formation and propagation were observed on this sample (and not on samples S2 or O2) both inside and outside the wear track (Fig. 10). This behaviour can be related to the higher penetration depth of the sliding ball on sample O3. Since the solid area fraction was rather low (because of the high porosity), applied load caused a very high local pressure, which could trigger crack formation relatively easily. Hence, the fretting tests were rather destructive for sample O3. Yet, despite the cracks, the majority of pores were durable. Actually, it is remarkable that despite of the wear and debris formation under the applied harsh reciprocating sliding conditions, the tubular structure was retained after the filling-up process for all samples. Therefore, the nanoporous AAO film can survive without cracks propagation under dry fretting conditions, which is of paramount importance for most practical applications. Nevertheless, it must be taken into account that the interplay between several parameters such as load, amplitude, frequency, AAO pore diameter and porosity degree is vital to form the self-lubricating and crack-preventing film on the AAO membranes that will protect tribological system against failure caused
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Fig. 10. Scanning electron microscopy (SEM) images of wear tracks after 500 cycles under a normal load of 1000 mN for S2, O2 and O3 samples.
by non-equilibrium behaviour.
between
abrasive
and
adhesive
wear
4. Conclusions In this study, anodic aluminium oxide membranes obtained from two different electrolytic baths were structurally characterized and tested for their mechanical and tribological behaviour. The anodization conditions were tuned in order to tailor the porosity degree and the pore diameter. Mechanical properties and frictional behaviour of the obtained AAO membranes were investigated and analysed as a function of the porosity percentage and the anodization bath. The main conclusions from this work can be summarized as follows: ▪ Normal loading by nanoindentation does not reveal any major crack formation on AAO. Penetration depth of the indenter is highly affected by the porosity percentage. Although no major cracks are formed, minor cracks are present on samples containing sulphur as contamination. Such an anion contamination presumably influences the deformation mechanism, particularly the elastic recovery. In particular, carbon from oxalic bath makes the surface more rigid. ▪ Hardness varies in accordance with the porosity level. Increment in the porosity percentage causes a decrease of hardness. ▪ The Young's modulus of AAO also depends on the porosity level and the anodizing bath. For a given porosity, oxalic acid anodizing produces an oxide layer with higher elastic modulus as compared to samples prepared from the sulphur-based electrolyte.
▪ AAO exhibits a very high coefficient of friction during reciprocating sliding against corundum. Wear causes the formation of a thick, compacted and sparsely distributed debris layer. This top layer smear out over the wear track and induces a high adhesion between AAO and corundum counterbody. Pores are partially filled-up with debris at the beginning of the reciprocating sliding tests. Increment in the pore size promotes the entrapment of fine debris. Interestingly, the porous structure is retained on all samples underneath the compacted debris layer. Acknowledgements This research was funded by a FP7 grants: NANOALLOY (909407/ 252407) and Oil & Sugar (295202). Partial financial support from the Catalan 2014-SGR-1015, the Spanish MINECO (MAT201127380-C02-01) and Moldovan National (72/ind) Foundations are also acknowledged. References [1] M. Almasi Kashi, A. Ramazani, M. Raoufi, A. Karimzadeh, Thin Solid Films 518 (2010) 6767. [2] T.-H. Fang, T.H. Wang, C.-H. Liu, L.-W. Ji, S.-H. Kang, Nanoscale Res. Lett. 2 (2007) 410. [3] S.F. Hulbert, in: L.L. Hench, J. Wilson (Eds.), An Introduction to Bioceramics, World Scientific Publishing Co. Pte. Ltd, Singapore, 1993. [4] P. Christel, A. Meunier, J.-M. Dorlot, J.-M. Crolet, J. Witroet, L. Sedel, P. Boutin, in: P. Ducheyne, J.E. Lemons (Eds.), Bioceramics: Material Characteristics Versus In Vivo Behaviour, New York Academy of Science, New York, 1988. [5] M. Karlsson, PhD Thesis, Uppsala, Finland, 2004. , S.K. Youn, E. Pellicer, S. Schuerle, J. Sort, S. Fusco, H. Gyu [6] M.A. Zeeshan, S. Pane Park, B.J. Nelson, Adv. Funct. Mater. 23 (2013) 823e831.
N. Tsyntsaru et al. / Materials Chemistry and Physics 148 (2014) 887e895 [7] J.-P. Salvetat, G. Andrew, D. Briggs, J.-M. Bonard, R.R. Bacsa, A.J. Kulik, €ckli, N.A. Burnham, L. Forro , Phys. Rev. Lett. 82 (1999) 944. T. Sto [8] Y. Wang, L. Xia, J. Ding, N. Yuan, Y. Zhu, Trib. Lett. 49 (2013) 431. [9] M. Kylan McQuaig Jr., A. Toro, W. Van Geertruyden, W.Z. Misiolek, J. Mater. Sci. 46 (2011) 243. [10] J.P. Tu, C.X. Jiang, S.Y. Guo, X.B. Zhao, M.F. Fu, Wear 259 (2005) 759. [11] A.K. Kothari, E. Konca, W.S. Brian, K. Jian, H. Li, Z. Xia, W. Ni, R. Hurt, J. Mater. Sci. 44 (2009) 6020. [12] Z. Xia, L. Riester, B.W. Sheldon, W.A. Curtin, J. Liang, A. Yin, J.M. Xu, Rev. Adv. Mater. Sci. 6 (2004) 131. [13] S. Ko, D. Lee, S. Jee, H. Park, K. Lee, W. Hwang, Glass Phys. Chem. 31 (2005) 356. [14] K.Y. Ng, Y. Lin, A.H.W. Ngan, Acta Mater. 57 (2009) 2710. [15] T.H. Fang, T.H. Wang, C.H. Liu, L.W. Ji, S.H. Kang, Nanoscale Res. Lett. 2 (2007) 410. [16] A.P. Samantilleke, J.O. Carneiro, S. Azevedo, V. Teixeira, T. Thuy, J. Nano Res. 25 (2013) 77. [17] P. Skeldon, H.W. Wang, G.E. Thompson, Wear 206 (1997) 187. [18] C.X. Jiang, J.P. Tu, S.Y. Guo, M.F. Fu, X.B. Zhao, Acta Metall. Sin. (Engl. Lett.) 18 (2005) 249. [19] N. Hu, S. Ge, L. Fang, J. Cent. South Univ. Technol. 18 (2011) 1004. [20] G.D. Sulka, S. Stroobants, V. Moschalkov, G. Borghs, J.-P. Celis, J. Electrochem. Soc. 149 (2002) D97. lu, MSc. Thesis, Istanbul, Turkey, 2011. [21] I. Pas¸aog [22] W.C. Oliver, G.M. Pharr, J. Mater. Res. 19 (2004) 3.
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[23] H. Mohrbacher, J.-P. Celis, J.R. Roos, Trib. Int. 28 (1995) 296. [24] D. Drees, J.-P. Celis, S. Achanta, Surf. Coat. Technol. 188e189 (2004) 511. [25] H.S. Kim, D.H. Kim, W. Lee, S.J. Cho, J.H. Hahn, H.S. Ahn, Surf. Coat. Technol. 205 (2010) 1431. s, J. Ferre -borrull, L.F. Marsal, J.-P. Celıs, Surf. [26] L. Vojkuvka, A. Santos, J. Pallare Coat. Technol. 206 (2012) 2115. [27] Y.C. Wang, I.C. Leu, M.H. Hon, J. Appl. Phys. 95 (2004) 1444. [28] A.C. Fischer-Cripps, in: F.F. Ling (Ed.), Nanoindentation, Springer, New York, 2004. [29] http://www.memsnet.org/material/aluminumoxideal2o3bulk/. [30] M. Asmani, C. Kermel, A. Leriche, M. Ourak, J. Eur. Ceram. Soc. 21 (2001) 1081. re, A. Vincent, Y. Bre chet, J. Appl. Phys. 80 (1996) 3772. [31] D. Bellet, P. Lamagne , V. Panagiotopoulou, S. Fusco, K.M. Sivaraman, S. Surin ~ ach, [32] E. Pellicer, S. Pane , B.J. Nelson, J. Sort, Int. J. Electrochem. Sci. 7 (2012) 4014. M.D. Baro [33] R.E. Williford, X.S. Li, R.S. Addleman, G.E. Fryxell, S. Baskaran, J.C. Birnbaum, C. Coyle, T.S. Zemanian, C. Wang, A.R. Courtney, Microporous Mesoporous Mater. 85 (2005) 260. [34] S. Ko, D. Lee, S. Jee, H. Park, K. Lee, W. Hwang, Thin Solid Films 515 (2006) 1932. [35] F. Tancret, F. Osterstock, Philos. Mag. 83 (2003) 125. [36] S. Carioua, F.-J. Ulma, L. Dormieux, J. Mech. Phys. Solids 56 (2008) 924. [37] J. Biener, A.M. Hodge, A.V. Hamza, L.M. Hsiung, J.H. Satcher, J. Appl. Phys. 97 (2005) 024301. [38] N. Baizura, A.K. Yahya, J. Non-Cryst. Solids 357 (2011) 2810.
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Mechanical and frictional behaviour of nano-porous anodised aluminium N. Tsyntsaru a, *, B. Kavas b, c, J. Sort d, M. Urgen b, J.-P. Celis e a
Institute of Applied Physics of ASM, 5 Academy str., Chisinau, MD 2028, Moldova Istanbul Technical University, Department of Metallurgical and Materials Engineering, 34469 Maslak, Turkey c Ford Otomotiv San A.S., Istanbul, Turkey d Catalana de Recerca i Estudis Avançats (ICREA) and Departament de Física, Universitat Auto noma de Barcelona, E-08193 Bellaterra, Spain Institucio e KU Leuven, Dept. MTM, Kasteelpark Arenberg 44, B-3001, Belgium b
h i g h l i g h t s
g r a p h i c a l a b s t r a c t
Well-ordered porous AAO with pore diameters between 16 and 75 nm were produced. Porosity and composition of electrolytic baths influence the mechanical properties. Ball-on-flat configuration was used in tribological testing under dry conditions. Adherent tribolayer consisting of fine worn particles prevents AAO from cracking. Testing parameters are moreover essential to envisage AAO practical applications.
a r t i c l e i n f o
a b s t r a c t
Article history: Received 13 March 2014 Received in revised form 28 August 2014 Accepted 30 August 2014 Available online 17 September 2014
The porous structure of anodic aluminium oxide (AAO) can be used in versatile applications such as a lubricant reservoir in self-lubricating systems. Such systems are subjected to biaxial loading, which can induce crack formation and propagation, ultimately leading to catastrophic mechanical failure. In this study, the mechanical and tribological behaviour of AAO, prepared from two different types of electrolytes (sulphuric and oxalic acids), are studied in detail. The electrolytic conditions are adjusted to render highly tuneable average pore diameters (between 16 and 75 nm), with porosity levels ranging from 9% to 65%. Well-ordered porous AAO are produced by two-step anodization at rather low temperatures. Mechanical properties, mainly hardness and Young's modulus, are investigated using nanoindentation. Both the porosity degree and the composition of the electrolytic baths used to prepare the AAO have an influence on the mechanical properties. Ball-on-flat configuration was used to estimate the tribological behaviour under dry conditions. No major cracks were observed by scanning electron microscopy, neither after indentation or fretting tests. In the running-in period of tribology experiments the pores
Keywords: Nanostructures Oxides Coatings Mechanical properties Tribology and wear
* Corresponding author. E-mail addresses: [email protected], [email protected] (N. Tsyntsaru), [email protected] (B. Kavas), [email protected] (J. Sort), [email protected] (M. Urgen), [email protected] (J.-P. Celis). http://dx.doi.org/10.1016/j.matchemphys.2014.08.066 0254-0584/© 2014 Elsevier B.V. All rights reserved.
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were filled with debris. This was followed by the formation of a highly adherent tribolayer e a third body consisting of fine worn particles originated from both the sample and the counterbody. Pore diameter and porosity percentage are found to strongly affect hardness and Young's modulus, but they do not have a major effect on the frictional behaviour. © 2014 Elsevier B.V. All rights reserved.
1. Introduction The unprecedented growth in different areas of nanotechnology has prompted the rapid development of novel routes (advanced lithography procedures, wet chemistry, self-assembly, etc.) to prepare materials with reduced lateral dimensions. The characterization of anodic aluminium oxide (AAO) is of outmost importance since this material has possible applications in fields such as catalysis, chemical sensors, biosensors, filters, humidity sensors, as well as ceramic materials for implant materials [1e5]. The AAO can also serve as templates for the growth of electrodeposited nanowires [6] and as a support for measuring the mechanical properties of nanocarbon tube ropes [7], amongst others. AAO is also potentially interesting for tribological applications since the nanoporous structure can be used as a reservoir or template for solid lubricants (e.g., nano-tubes or nano-fibres) to form self-lubricating structures [8]. AAO nanostructures are typically composed of two regions: a relatively pure alumina inner layer, consisting entirely of Al2O3, and an acid anion-contaminated outer layer, resulting from the incorporation of anions into the alumina structure during anodizing [9]. The presence of these two regions together with the occurrence of different crystallographic phases (e.g., amorphous, a-Al2O3 and gAl2O3, that stem from different heat treatment conditions) affect the mechanical strength. Friction and wear performance of filled-in AAO have been studied to some extent [10]. Fundamentally, homogenous and heterogeneous (i.e., composite) ceramics behave in different ways when subjected to normal loading. Namely, homogeneous ceramics typically respond to the intense confined shear under the sharp indenter tip by an elastic deformation followed by the formation of radial cracks due to elastic mismatch. Conversely, heterogeneous ceramics can accommodate elastic strains much more efficiently, hence minimizing crack propagation and showing improved fracture toughness [11]. In spite of the previous studies on the mechanical properties of AAO [12e16], there is still lack of knowledge in some particular issues, such as the detailed understanding on how the porous structure of AAO and the concomitant mechanical properties (hardness, Young's modulus) evolve during mechanical deformation. Information concerning the tribological properties of AAO templates is also rather limited. Some literature exists on the tribological characterization and friction behaviour of AAO filled up with solid or liquid lubricants [17e19]. However, the frictional behaviour of AAO itself has been poorly investigated yet. In this study, we investigate (i) the mechanical properties of AAO using nanoindentation with a Berkovich indenter tip, and (ii)
the tribological response in a ball-on-flat sliding tester. The effect of porosity level, pore diameter, and the type of anodizing electrolyte on mechanical and tribological properties are described. Particular emphasis is given on how the mechanical tests modify porosity and, consequently, on how they affect the mechanical properties of AAO. The acquired understanding of the evolution of the physical response and the wear behaviour of AAO under mechanical loading is of crucial importance in many mechanical systems, and triggers new ideas for innovative application areas. 2. Experimental 2.1. Surface preparations and anodizing Al sheets with a thickness of 0.25 mm (99.99%, Alfa Aesar Johnson Matthey GmbH) were electropolished at a constant current density of 500 mA cm1 for 1 min in an electrolyte consisting of HClO4 (60 wt. %) and C2H5OH (abs.) in a volumetric ratio of 1:4. Temperature was kept constant at ~10 C. Samples were examined using scanning-electron microscopy (SEM) and non-contact white light interferometer (WYKO) before and after electropolishing for the estimation of Ra (arithmetical mean roughness) and Rz (tenpoint mean roughness) values. Surface roughness decreased after electropolishing (i.e. Ra from 740 nm to 250 nm, and Rz from ~7 mm to ~3.2 mm). The anodizing of electropolished samples was carried out in two different electrolytes, namely 20 wt.% H2SO4 [20] and 0.3 M H2C2O4 [21], in order to produce AAO with different characteristics (e.g., pore size or interpore distance). The synthesis procedures used are described in Table 1. The samples were anodised for various electrolysis durations, so as to obtain AAO with an average thickness of ~17 mm. Two-step anodisation was successful, as can be seen from the SEM images (Fig. 1). Well-ordered, hexagonally distributed pore arrays were obtained with different pore diameters and porosity degree (Table 2). In order to simplify the nomenclature of the samples, samples produced in sulphuric acid are denoted hereafter as S1 and S2, respectively, while those anodized in oxalic acid and pore widened afterwards are identified as O1, O2 and O3 (Tables 1 and 2). Pore widening was performed at room temperature (~25 C) by immersion of anodized in oxalic acid AAO in H3PO4 (5 wt.%) for different durations namely 32 min (O2) or 75 min (O3) to obtain variable pore diameters. Similar pore diameters were produced for samples O1 and S2 in order to assess the influence of the type of electrolytic bath on the wall thickness and the mechanical
Table 1 Overview of the samples preparation conditions. Sample name
Anodizing, I step
S1 S2 O1 O2 O3
15 V Sulphuric 500 rpm, 1 C, 10 min 21 V Oxalic 40 V, 800 rpm, 25 C, 60 min
Oxide removal 10 min 15 min 60 min
CrO3 (1.8 wt. %) þ H3PO4 (6 wt. %) 60 C
Anodizing, II step
Pore widening
15 V, 97 min Sulphuric 500 rpm, 1 C 21 V, 20 min Oxalic 70 V, 800 rpm, 5 C, 60 min
No No 32 min 75 min
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properties. Chemical composition of AAO was quantified by energydispersive spectroscopy (EDS). The EDS analyses revealed the presence of carbon and sulphur contaminations in the AAO samples produced from oxalic and sulphuric acids baths, respectively. 2.2. Mechanical and tribological characterization To study the mechanical behaviour of anodized aluminium oxides, nanoindentation tests were performed at normal loads varying from 5 to 100 mN using an indenter from CSM Instruments, equipped with a Berkovich-type tip. Loading rates were chosen to be twice the value of the peak applied load per minute. Indentation imprints were observed by SEM. Loading vs. penetration depth data were recorded. Hardness values were evaluated using the method of Oliver & Pharr [22] from a total of 5 indentation curves. Wear behaviour was investigated in relation to pore diameter, porosity percentage and oxide composition. Tribological tests were performed in the KUL-MTM fretting mode I apparatus [23] and a Modular Universal Surface Tester (MUST, Falex Tribology N.V., Belgium) in a ball-on-flat configuration operated under reciprocating mode. The counterbody was a 5 mm diameter corundum ball reciprocating on AAO samples at a frequency of 1 Hz, a displacement of 100 mm, under normal loads of 40 mN, 80 mN (mesoscale) and 1000 mN (macroscale). The mesoscale and the macroscale terms are derived from the displacement and normal loads used as is given e.g. in Ref. [24]. The test temperature and humidity were kept constant at 23 C and 50% RH, respectively. It should be noted that steel balls had been used as counterbody in previous works [25], and this resulted in tribochemical reactions during sliding. Here, to avoid these effects, corundum balls were used as counterbody, i.e., the same material as the base membranes (AAO). Samples were degreased before the reciprocating sliding tests. After sliding tests, the samples were ultrasonically cleaned in ethanol for 7e10 min to remove the debris. 3. Results and discussion 3.1. Mechanical characterization of nanoporous AAO by nanoindentation
Fig. 1. Scanning electron microscopy (SEM) images of produced anodised aluminium oxides (AAO).
SEM images of the nanoindentation imprints are shown in Fig. 2. Since AAO is a ceramic (i.e., non-ductile material), crack formation and propagation could be a priori expected as a result of the indentation compressive and shear stresses. However, SEM images revealed that the investigated AAO did not show major cracks either inside or near the imprint, as opposed to some previous studies [14]. The lack of major cracks might be related to the intrinsic structure of the investigated AAO or to the difference in strain rate and other test parameters, as compared to the studies carried out by other authors. Similar results with our tests were obtain in [26], but with less evidence of different behaviour between samples anodised in sulphuric and oxalic acids. The anodized layers show differences in structure when compared to bulk brittle Al2O3. They contain substantial amounts of OH groups, as was previously evidenced [27]. These OH groups cause an increase of the overall intrinsic plasticity, thus inhibiting crack nucleation. The high-magnification images (see Fig. 2, insets) reveal that minor cracks are actually formed within the imprinted area during the course of the indentation experiments in samples S1 and S2. Sulphur content and the structure of AAO produced in sulphuric acid might have made the oxide more brittle. On the contrary, samples O1, O2 and O3 do not show any sign of cracking. For given indentation conditions (i.e., maximum load) the number of cracks formed as a result of indentation was less on sample S2, which contained wider pores than sample S1. This suggests that the
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Table 2 Overview of the morphological features of the investigated samples. Sample
Pore diameter, nm
Interpore distance, nm
Porosity, %
S1 S2 O1 O2 O3
~16 ~27 ~27 ~45 ~75
~39 ~56 ~94 ~94 ~94
~15 ~20 ~9 ~20 ~65
pore diameter affects the formation of cracks. When comparing O1 and S2, which had both approximately the same pore diameters, we did not observe any cracks on O1. This may be related to the differences between the structural composition of AAO films produced in sulphuric and oxalic acids. No cracks were observed either on samples O2 or O3. Loadepenetration depth curves were measured on all samples. As an example, Fig. 3 shows representative indentation curves obtained on sample O3 for maximum applied loads of 5, 20, 50 and 100 mN. As it was noticed on all loadingeunloading curves, no popin events corresponding to major cracks or phase transformations took place during the nanoindentation experiments. The indentation curves were rather smooth without any obvious sudden changes. The average values of maximum and final penetration depths, as well as the elastic recovery percentages, were estimated from the indentation curves (see Table 3). The results indicate that the elastic recovery of AAO produced in sulphuric acid, regardless of their pore diameter or porosity degree, is larger than for samples produced in oxalic acid. Indeed, although O2 and S2 show similar porosity percentages (Table 2), a significant difference in the elastic recovery is observed (i.e., while O2 has a recovery of ~24%, that of S2 is ~41%). Similarly, although the pore diameters of samples O1 and S2 are similar, a clear difference is again observed on the elastic recovery. Remarkably, samples produced in oxalic acid electrolyte show similar recovery percentages in spite of their different pore diameter and porosity level. This seems to suggest that, rather than the porosity degree and pore size, the type of electrolytic bath used for anodizing plays a major role in the elastic recovery of the samples. Contrary to what is observed for the elastic recovery, the porosity level is found to have a strong influence on the hardness and reduced Young's modulus. Namely, the maximum penetration depth attained in sample O3 (with a porosity of ~65%) is much larger than on samples O2 (with porosity of ~20%) or O1 (with porosity of ~9%). As a consequence, sample O3 is mechanically softer (see Fig. 4). The effect of porosity on the Young's modulus of the AAO porous films was also analysed. As shown in Fig. 5, there is a clear trend for Er to decrease with the increase of porosity level. This is particularly clear for the oxalic samples, especially at low indentation loads, where Er decreases from 110 GPa (for 9% porosity) to 103 GPa (for 20% porosity) and to 60 GPa (for 65% porosity). The reduced Young's modulus takes into account the elastic displacements that occur in both the specimen, with Young's modulus E and Poisson's ratio n, and the diamond indenter, with elastic constants Ei ¼ 1140 GPa and ni ¼ 0.07. The relationship between Er and E can be expressed as [28]:
1 1 v2 1 v2i þ ¼ Er E Ei
(1)
For bulk Al2O3, the Poisson's ratio is around 0.24 [29]. This value could be slightly lower if one takes into consideration the existence of porosity [30]. Assuming that n z 0.24, the Young's modulus of the porous Al2O3 would be: E z 114 GPa for 9% porosity, 106.6 GPa
Fig. 2. Scanning electron microscopy (SEM) images of the AAO templates indented using a maximum load of 100 mN. Shown in the insets are high magnification images of the indented regions.
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Fig. 3. Loadingeunloading nanoindentation curves corresponding to sample O3 for maximum applied loads of 5, 20, 50 and 100 mN.
for 20% porosity and 59.7 GPa for 65% porosity, hence in practice very similar to Er. The porosity level is known to have a strong influence on the elastic constants of metallic and ceramic materials. In a first approximation, it has been shown that [31e33]:
Eporous ¼ Ebulk
rporous n rbulk
(2)
where rporous/rbulk is the relative density. The exponent n has been reported to be equal to 2 for open-celled foams or sponge-like structures [31,32], whereas n ¼ 3 is generally assumed for materials whose porous structure is arranged forming a honeycomb hexagonal array normal to the surface [33]. In the case of porous Al2O3, values close to n ¼ 1 have been reported [33]. The relative density is related to the porosity volume fraction:
rporous ¼1P rbulk
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Fig. 4. Dependence of hardness on the maximum penetration depth for the investigated AAO samples.
Concerning the hardness tests (Fig. 4), the values obtained for samples O1, O2 and O3 are also lower than those reported for nonporous anodized alumina, which have been reported to be around 13 GPa [34]. The decrease of hardness (or compressive yield stress) with porosity is also a well-documented effect [35e37] and has been modelled using finite element simulations of nanoindentation curves [36]. An equation analogous to Eq. (2) may be used to correlate the yield stress of the porous structure with that of the bulk solid material:
sporous ¼ C2 sbulk
rporous rbulk
m (4)
where C2 ¼ 0.3 and m is an exponent [37]. Although the relationship between hardness and yield stress in bulk oxide ceramic is
(3)
The Young's modulus of bulk (non-porous) amorphous anodized Al2O3 (non-heat treated) has been reported to be around 147 GPa [34]. Using Equation (2), values of n equal to 2.8, 1.5 and 0.85 are obtained for indentations made using a maximum load of 5 mN, for samples O1, O2 and O3 with 9%, 20% and 65% porosity, respectively. These values of n fall within the range of expected values for porous materials. The values of n remain similar for samples O1 and O2, regardless of the applied indentation load. Conversely, for sample O3, the exponent n increases with the applied load, up to n ¼ 1.2, probably due to the densification effects observed during nanoindentation (see SEM image in Fig. 2).
Table 3 Summary of the values of maximum depth, final depth and elastic recovery for the different AAO templates, after indentation tests using a maximum load of 100 mN. Sample
Max. penetration depth, nm
Final depth, nm
Recovery %
S1 S2 O1 O2 O3
~1320 ~1325 ~1010 ~1110 ~1650
~840 ~780 ~710 ~845 ~1260
~36 ~41 ~29 ~24 ~24
Fig. 5. Dependence of the reduced Young's modulus on the maximum penetration depth for the different investigated AAO membranes.
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often taken as Hbulk z 1.6sbulk [38], a quantitative relationship between hardness and yield stress in porous materials has not been readily established yet. Assuming that Hporous ¼ 1.6sporous [38], Eq. (4) would give m z 0.4 for the samples with 65% porosity and applied load of 5 mN (and lower values for the samples with lower porosity degree), which are much smaller than the exponent obtained in an open-cell metallic foam (m ¼ 1.5) [37], but in agreement with the decrease of hardness with porosity observed by other authors on anodized alumina [34]. The samples S1 and S2 show relatively similar hardness values at all applied loads, varying between ~4.5 and ~6 GPa, except at small loads where there is large scatter (see Fig. 4). The Young's modulus is also similar for these two samples (Fig. 5). This is probably due to the very small difference in the porosity percentage between these two samples. The unexpected larger hardness recorded on sample S1 at 5 mN could be linked to the rougher surface of this sample. 3.2. Frictional behaviour of nanoporous AAO at meso- and macroscale loading Nanoindentation tests provide evidence of the range of normal forces that AAO templates can withstand without being catastrophically damaged. In analogy, the evaluation of the tribological behaviour is important to assess the range of forces that can be applied under sliding to this material before it undergoes dramatic frictional effects. In order to evaluate how the porous structure behaves when subjected to bidirectional loading, two types of reciprocating sliding cycles were set: (i) 50 cycles to reveal the initial evolution of wear mechanism and (ii) 500 cycles to observe the frictional behaviour already in the steady-state regime. Wear tracks were examined using SEM after 50 cycles performed at relatively low loads (40 and 80 mN) (Fig. 6). The observations revealed an interesting response; namely, pores started to be filled-up with fine debris particles without any cracking. This behaviour proved that the porous tubes were strong enough to resist bidirectional loading at this level of applied forces. Tangential force versus displacement hysteresis loops showed that the tests were carried out under gross slip regime. In the running-in period of meso scale tests (i.e., at applied load of 40 and 80 mN), the tangential force was first rather low and increased with the number of cycles (Fig. 7). This demonstrates that abrasive wear at the beginning causes the formation of wear particles which act as third body making the relative motion more difficult due to adhesion. As a result, the wear mechanism is a combination of abrasive and adhesive wear, hence resulting in a high coefficient of friction (Fig. 7). At the steady state, the coefficient of friction reaches approximately 0.9. For higher
Fig. 7. Evolution of the coefficient of friction for S2 with the number of cycles under different applied loads.
loads (i.e., macro-scale loading at 1000 mN), the running-in period is followed with a decrease of the coefficient of friction until a value close to 1 is reached in the steady-state (Fig. 7), quite close to that measured under meso-scale loading. Increasing normal load from 40 mN (Fig. 8a) to 80 mN (Fig. 8b) under the same displacement amplitude brings about a decrease in effective sliding distance in the contact with sliding cycles. This can also be linked to the adhesion of debris or the penetration of the ball into the wear track. As at macro scale loading (Fig. 8c), tangential force shows a decrease with the number of reciprocating sliding cycles. The difference between the evolution of hysteresis loops under 1000 mN and meso-scale loading can be linked to the penetration depth of the counterbody which was deeper at higher normal loads. Once the ball creates a wear track by abrasive wear, the relative motion becomes easier, revealed by a decreasing tangential force. The stiffness of the macro-scale test machine is larger than that of meso-scale tester [26]. Thus, the obvious effect of adhesive wear observed at meso-scale loads is less apparent at macro-scale loads. Evolution of the coefficient of friction with the number of cycles for all samples, at 1000 mN normal load, is given in Fig. 9. A high coefficient of friction was expected for all samples since a corundum counterbody was used (i.e., the same material as the test
Fig. 6. Scanning electron microscopy (SEM) images of S2 samples after 50 cycles fretting under 40 mN (a) and 80 mN (b).
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Fig. 8. Fretting hysteresis loops registered for S2 under 40 mN (a), 80 mN (b) and 1000 mN.
sample). Similar coefficient of friction dependencies were noticed on samples O2, S1 and S2 (which had rather similar porosity percentages). The coefficient of friction recorded on these samples was ~1.35 at the beginning and decreased down to ~1.1 once steady state sliding is established. The resemblance of the results can indeed be expected from the similar porosity percentages, which allow analogous amounts of debris entrapped in the pores and also result in similar penetration depths. The coefficient of friction recorded on O1 is higher than on S1, S2 and O2 probably due to a lower porosity which means bigger amount of aluminium oxide in the contact area. In this sense, sample O3, with the highest porosity level, should exhibit the lowest coefficient of friction, which was not noticed during our experiments. The reason for the relatively high coefficient of friction (~1.5) in this highly porous sample could be the high penetration depth achieved on O3 (similar to what was observed during nanoindentation tests). As the ball goes deeper inside the oxide layer, it becomes surrounded by a larger amount of nanotubes as if it became buried within the AAO. As a result, higher tangential force is required to carry out the reciprocating motion. Furthermore, this sample is also able to entrap larger amounts of debris particles and retards the formation of a top-layer consisting of debris that induces a decrease in the coefficient of friction. SEM images of wear tracks produced on S2 sample after 500 cycles are shown in the upper row of Fig. 10. Debris was sparsely
Fig. 9. Evolution of the coefficient of friction of all samples under an applied load of 1000 mN.
distributed on the surface in piles forming a third-body tribolayer. The tribolayer consisting of compacted wear particles exhibits an encrusted structure. The debris adhering locally at some parts in the wear track causes a complex wear mechanism stemming from a combination of abrasive and adhesive wear together with a high coefficient of friction (Fig. 9). This debris is highly adherent and remains visible even after ultrasonically cleaning the samples for 10 min in ethanol. Wear tracks on sample O2 were similar to those produced on S2 (Fig. 10). Compacted debris particles inside the wear track and a layered structure with sparsely distributed debris were observed. Indeed, analogous complex wear mechanism as that on S2 is expected on O2, thus resulting in a high coefficient of friction. On sample O2 pores partially filled with debris in areas where compacted debris did not adhere, were still visible after 500 cycles. Conversely, the wear tracks on the highly porous sample O3 show some different features (Fig. 10). The light greyish areas observed on samples O2 and S2 at low magnification were not present on sample O3. Moreover, at higher magnification, the layered and compacted debris were neither as dense nor as thick as those seen on samples S2 and O2. Pores remained more visible. Because of the larger pore size, the nanotubes were easily filled with most of the debris. Such progressive filling probably inhibits, to some extent, the formation of a dense and thick top layer. Nevertheless, as can be envisaged from the wear tracks, the nanotubes were just overfilled after 500 cycles and in case of running the reciprocating sliding tests for longer durations, similar top layers as on samples S2 and O2 would probably form. Although the tubular structure of sample O3 was not significantly damaged, crack formation and propagation were observed on this sample (and not on samples S2 or O2) both inside and outside the wear track (Fig. 10). This behaviour can be related to the higher penetration depth of the sliding ball on sample O3. Since the solid area fraction was rather low (because of the high porosity), applied load caused a very high local pressure, which could trigger crack formation relatively easily. Hence, the fretting tests were rather destructive for sample O3. Yet, despite the cracks, the majority of pores were durable. Actually, it is remarkable that despite of the wear and debris formation under the applied harsh reciprocating sliding conditions, the tubular structure was retained after the filling-up process for all samples. Therefore, the nanoporous AAO film can survive without cracks propagation under dry fretting conditions, which is of paramount importance for most practical applications. Nevertheless, it must be taken into account that the interplay between several parameters such as load, amplitude, frequency, AAO pore diameter and porosity degree is vital to form the self-lubricating and crack-preventing film on the AAO membranes that will protect tribological system against failure caused
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Fig. 10. Scanning electron microscopy (SEM) images of wear tracks after 500 cycles under a normal load of 1000 mN for S2, O2 and O3 samples.
by non-equilibrium behaviour.
between
abrasive
and
adhesive
wear
4. Conclusions In this study, anodic aluminium oxide membranes obtained from two different electrolytic baths were structurally characterized and tested for their mechanical and tribological behaviour. The anodization conditions were tuned in order to tailor the porosity degree and the pore diameter. Mechanical properties and frictional behaviour of the obtained AAO membranes were investigated and analysed as a function of the porosity percentage and the anodization bath. The main conclusions from this work can be summarized as follows: ▪ Normal loading by nanoindentation does not reveal any major crack formation on AAO. Penetration depth of the indenter is highly affected by the porosity percentage. Although no major cracks are formed, minor cracks are present on samples containing sulphur as contamination. Such an anion contamination presumably influences the deformation mechanism, particularly the elastic recovery. In particular, carbon from oxalic bath makes the surface more rigid. ▪ Hardness varies in accordance with the porosity level. Increment in the porosity percentage causes a decrease of hardness. ▪ The Young's modulus of AAO also depends on the porosity level and the anodizing bath. For a given porosity, oxalic acid anodizing produces an oxide layer with higher elastic modulus as compared to samples prepared from the sulphur-based electrolyte.
▪ AAO exhibits a very high coefficient of friction during reciprocating sliding against corundum. Wear causes the formation of a thick, compacted and sparsely distributed debris layer. This top layer smear out over the wear track and induces a high adhesion between AAO and corundum counterbody. Pores are partially filled-up with debris at the beginning of the reciprocating sliding tests. Increment in the pore size promotes the entrapment of fine debris. Interestingly, the porous structure is retained on all samples underneath the compacted debris layer. Acknowledgements This research was funded by a FP7 grants: NANOALLOY (909407/ 252407) and Oil & Sugar (295202). Partial financial support from the Catalan 2014-SGR-1015, the Spanish MINECO (MAT201127380-C02-01) and Moldovan National (72/ind) Foundations are also acknowledged. References [1] M. Almasi Kashi, A. Ramazani, M. Raoufi, A. Karimzadeh, Thin Solid Films 518 (2010) 6767. [2] T.-H. Fang, T.H. Wang, C.-H. Liu, L.-W. Ji, S.-H. Kang, Nanoscale Res. Lett. 2 (2007) 410. [3] S.F. Hulbert, in: L.L. Hench, J. Wilson (Eds.), An Introduction to Bioceramics, World Scientific Publishing Co. Pte. Ltd, Singapore, 1993. [4] P. Christel, A. Meunier, J.-M. Dorlot, J.-M. Crolet, J. Witroet, L. Sedel, P. Boutin, in: P. Ducheyne, J.E. Lemons (Eds.), Bioceramics: Material Characteristics Versus In Vivo Behaviour, New York Academy of Science, New York, 1988. [5] M. Karlsson, PhD Thesis, Uppsala, Finland, 2004. , S.K. Youn, E. Pellicer, S. Schuerle, J. Sort, S. Fusco, H. Gyu [6] M.A. Zeeshan, S. Pane Park, B.J. Nelson, Adv. Funct. Mater. 23 (2013) 823e831.
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