Thin Solid Films 536 (2013) 179–186
Contents lists available at SciVerse ScienceDirect
Thin Solid Films
journal homepage: www.elsevier.com/locate/tsf
The structure and properties of thin aluminum coatings
Alexander L. Volynskii a, Daria A. Panchuk a, Sergey L. Bazhenov b, Mikhail Yu. Yablokov b, Alla B. Gilman b,
Anastasia V. Bolshakova a,⁎, Larisa M. Yarysheva a, Nikolai F. Bakeev b
a
b
Chemistry Department, Moscow State University, 119991, Moscow, Russia
Institute of Synthetic Polymeric Materials, Profsoyuznaya Street 70, 117393 Moscow, Russia
a r t i c l e
i n f o
Article history:
Received 28 December 2011
Received in revised form 15 March 2013
Accepted 21 March 2013
Available online 2 April 2013
a b s t r a c t
Mechanical properties and structure of a thin aluminum coating were investigated. Experimental methods of
measuring yield stress, strength and fracture strain of ultrathin coatings were developed. Yield stress,
strength and fracture strain of aluminum increase with decrease in the coating thickness. The increase of
yield stress and strength was explained by reduction of the crystal size and by strain-hardening of metal.
© 2013 Elsevier B.V. All rights reserved.
Keywords:
Strength
Mechanical properties
Coating
Fragmentation
Buckling
1. Introduction
Thin films are increasingly used in coatings, optical reflectors etc.
[1]. Although the mechanical properties of submicrometer-thick
films are critical for their use, there are currently few methods for
measuring thin film properties. Conventional mechanical testing
devices are not sensitive enough to measure the forces resulting
from straining of submicrometer-thick films.
Scratch test [2–5] and indentation test [5–7] are sometimes used
to characterize the properties of thin rigid films on a substrate. In
the case of scratch test, scratches are made on the coated sample by
a diamond indenter, which is drawn across the surface. By indentation test the Young's modulus E and hardness of thin metal films
were determined. Both the hardness and the Young's modulus
depend on the film thickness [8,9]. Unfortunately, the scratch and
indentation tests do not allow to estimate the fracture strain and
strength.
In addition to indentation test, the Young's modulus and the
Poisson's ratio of metal coatings were determined by the X-ray
diffraction [10].
At scratch test, adhesion strength may be calculated from the size
of debonded zone and the load needed to remove the film from the
substrate. Although quantitative evaluation is easily realized in such
tests, the results are influenced by the strength of film and interface,
by film thickness [8,9], and residual stress [11]. Adhesion strength
was measured also by shear loading [12]. In the alternative approach
⁎ Corresponding author.
E-mail address: bolshakova@nanoscopy.ru (A.V. Bolshakova).
0040-6090/$ – see front matter © 2013 Elsevier B.V. All rights reserved.
http://dx.doi.org/10.1016/j.tsf.2013.03.050
adhesion is characterized by interface fracture toughness G [13,14].
Interface fracture toughness G represents the amount of energy needed to extend a crack by unit area along the substrate/coating plane. A
thin rigid coating supported by a soft substrate under tension load
can fracture into a number of parallel cracks oriented perpendicularly
to the loading direction. The crack in the coating may turn by 90° and
grow along the interface [13]. This process causes
debonding of the
pffiffiffiffiffiffiffiffiffiffiffiffiffiffi
coating. The coating debonds if the inequality 2EG=hbσ y is fulfilled,
where σy, E and h are yield stress, Young modulus and the thickness
of the coating respectively. If the coating thickness h exceeds some
critical value hc, the cracked coating debonds from the substrate. In
contrast, thin coatings do not debond from the substrate. The critical
thickness value hc for the Pt/polyethylene terephthalate (Pt/PET)
system was equal to 26 nm [13].
The Young's modulus of the coating or substrate was measured
also [15,16] using spontaneous buckling of a rigid coating on a soft
substrate under compressive or tensile load [17–21].
Mechanical, physical–chemical and other properties of thin films
depend on their thickness [22–28].
Thus, the properties of thin films may be measured by different
methods. However, there is a need for development of methods, especially for measurement of the coatings strength. The aim of this work
is to develop methods for measuring fracture strain, strength and
yield stress of thin aluminum coatings deposited on polymer films.
2. Experimental details
A commercial film of amorphous non-oriented polyethylene terephthalate (PET) was used as a substrate. The thickness of the film was
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100 μm. Samples, dumbbell in shape, were cut from the film. The
gauge of samples was 6 × 20 mm 2.
Thin aluminum coating was deposited on the polymer films by
thermal sputtering in vacuum in VUP-4 device. The thickness of the
coating was changed by variation of time of metal deposition. Control
measurements of the coating thickness were made by depositing an
aluminum coating on a glass plate. After that, the coating was
scratched with a sharp stick, which does not scratch the glass, and
the depth of the scratch was measured by an atomic-force microscope. Fig. 1 illustrates the method [29]. The light area in the left
part of the image is an aluminum layer, and the dark area on the
right is the glass surface. The coating thickness was proportional to
the time of metal deposition. The proportionality coefficient was
found from Fig. 1, thus determining the thickness of the aluminum
coating as a function of the deposition time.
Mechanical properties were studied with a standard “Instron-1122”
test machine. The samples were tested in tension at temperature 20
and 90 °C at a cross-head speed of 10 mm/min. After elongation of the
coated sample at temperature 90 °C (which is higher than the glass transition temperature Tg of PET substrate) it was cooled in the clamp of the
test machine to room temperature.
The samples after elongation were investigated in the scanning electron microscope (SEM) “Hitachi S-520”, working voltage is 20 kV,
transmission electron microscope (TEM) “LEO 912AB”, working voltage
is 80 kV, and atomic-force microscope “Nanoscope-IIIa” in the contact
and tapping modes, non-contact (fpN10S, length 100 μm, resonance
frequency 254 KHz), and contact (fpC10S, length 250 μm, force constant 0.1 N/m) cantilevers were used manufactured by State Research
Institute for Problems in Physics (Zelenograd, Russia). The average
values of the width of coating fragments and the period of the microwave were determined from SEM and scanning probe microscopy
(SPM) images using FemtoScan Online software [29].
3. Theoretical background
Fig. 2 shows SEM micrographs of a coated PET elongated at 90 °C
(2a) and at room temperature (2b). The sample 2a was elongated at
temperature higher than the glass transition temperature Tg (75 °C)
of PET substrate. The sample 2b was elongated at room temperature,
i.e. below Tg. The light bands, perpendicular to the elongation axis, are
the coating fragment bands. The dark bands show PET substrate made
visible due to opening of cracks in the coating. The coating in Fig. 2
is fragmented on several bands of approximately equal width. In
Fig. 2. SEM microphotographs of PET coated with 10 nm gold film and elongated by
50% at (a) 90 °C and (b) room temperature.
Fig. 1. SPM-image and a cross-section of a glass plate coated with scratched thin Al film.
addition, the coating is folded so that the surface is wavy, and the
wave crests are parallel to the tension direction. Spontaneous formation of a wave on an initially smooth surface is described in [17,18]. In
contrast, if a sample was elongated at temperature below Tg of the
polymer substrate (Fig. 2b), the wave-like folding do not appear. In
this case only the coating fragmentation is observed.
Both the coating folding and coating fragmentation are quite periodic. Both structures appear if the following conditions are fulfilled
[17,18]: 1) the coating is much thinner than the substrate and 2) the
elastic modulus of the coating is much higher than that of the substrate.
3) the coating is adhered to the substrate. The first two conditions are
fulfilled for a polymer with a metal coating at temperature higher
than the glass transition temperature of the polymer substrate. The
coating does not debond from the substrate if it is thin and its thickness
is lower than some critical value determined by the debonding fracture
toughness.
A.L. Volynskii et al. / Thin Solid Films 536 (2013) 179–186
181
3.1. Folding of the coating
The volume of a polymer at plastic deformation does not change,
and the film sample elongation is accompanied by contraction in
the transverse direction. Hence, the coating is tensioned along the
elongation axis and compressed in the transverse direction. As a
result, fragmentation of the coating and regular wave-like folding of
the coating are observed. Regular folding of the coating is shown in
Fig. 3 [17,18]. (Data from the atomic force microscope are already
shown in our previous publications [17,18]).
Assuming a sinusoidal waveform of the buckling instability and
elastic behavior of the coating and the substrate, the critical wavelength λ can be expressed by [18,20].
31=3
2
1−νs 2 Ec
5
λ ¼ 2πh4 1−νc 2 Es
Fig. 4. Stress distribution in a coating fragment adhered to polymer substrate.
where h is the coating thickness, ν is the Poisson's ratio, and E is the
Young's modulus. Subscripts c and s denote the film and substrate
respectively.
If ductile substrate and ductile coatings are used, the mechanism
of folding is plastic buckling instability. The coating in this case deforms as several plastic hinges. In this case the wave length is given
by equation:
λ ¼ 3h
σ yc
ð1 φÞ
σ ys
ð1Þ
where σyc and σys — yield stresses of the coating and the substrate respectively, φ = ΔL/Lo — the degree of the coating compression, ΔL
and Lo are reduction of the sample width and Lo is the initial sample
width. In the case of axial elongation, φ is the degree of the substrate
transverse contraction. The coating yield stress σyc may be determined if λ, h, σys and φ are measured.
3.2. Fragmentation of the coating
Regular fragmentation of the coating is observed both below and
above Tg of the substrate (Fig. 2). At low tensile strains, the fragmentation is caused by defects in the coating, and the width of fragment
bands is quite different [30–32]. The mechanism of fragmentation at
high strains changes, and the width of the coating fragments is
quite narrow [33]. Distribution of tensile stress in each fragment is
illustrated by Fig. 4. The stress at the ends of the fragment is equal
to zero. However, the stress increases along with the distance from
the fragment ends, reaching maximum values in the center of the
fragment. This leads to fracture of each fragment into two equal
parts. This is illustrated by Fig. 5. Fig. 5a shows the sample which
was elongated at high temperature. The stress in the substrate in
this case is relatively low. Several fragments with quite different
width are observed. Afterwards, the temperature was reduced to
room temperature, and the stress in the substrate increased. The sample was elongated again, and each fragment is divided into two equal
halves (Fig. 5b). Fragmentation of the coating into two equal parts
proceeds until polymer substrate is able to transfer a stress sufficient
for further fracture of each fragment. When this limiting value of
stress is reached, fragmentation stops. For this reason, several coating
fragments with rather narrow width are formed.
Fragment width distribution was analyzed in [30–33]. The average
width, L, of fragments is described by equation
L ¼ 3hσ =σ 0s
ð2Þ
where h is the coating thickness, σ* is the coating strength, and σos
is the stress in the substrate. Eq. (2) is analogous to Kelly equation describing ineffective fiber length in fiber-reinforced plastics [34]. The
coating strength σ* may be calculated if h, L and σo are measured.
3.3. Fracture strain of the coating
A method of measuring fracture deformation of metal coating was
developed in [35,36]. Fig. 6 demonstrates the coating fragments. Evidently, the sum of the coating fragments lengths in the elongation direction is equal to Lc = λcL0, where λc is the drawing ratio of the
coating and L0 — initial sample length. Similarly, the substrate length
is equal to Ls = λs L0. Then
L c λc 1 þ ε c
¼
¼
;
Ls λs 1 þ εs
where εc and εs are plastic deformations of the coating and the
substrate respectively, λc and λs are elongation ratios of the coating
and the substrate. Thus, the plastic strain of the coating is described
by equation:
εc ¼ Lc λs =Ls –1:
Fig. 3. Buckling of an elastic rod (a, b) and a coating on soft substrate (c, d).
ð3Þ
According to Eq. (3), the failure strain of the coatings may be
determined by measuring the coating fragments lengths Lc and the
total length of the sample Ls. Lc was measured from micrographs
(Fig. 6), and deformation of the substrate λt is given by sample elongation in the testing machine.
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Fig. 6. Schematic illustrating calculation of plastic tensile strain of coating from SEM data.
zero. Hence, Eq. (3) can be used to measure irreversible deformation
of the coatings which is equal to its fracture strain.
Relationships (1)–(3) constitute a basis of a simple and direct
method of estimating coatings properties. Strength, yield stress and
fracture strain of the coating were determined from the parameters
of coating fragmentation and folding, determined with microscopes.
4. Results and discussion
Fig. 5. SEM images of (a) PET film coated with 4 nm-thick platinum film elongated by
100% at 100 °C at a rate of 0.1 mm/min and (b) the same sample after additional
elongation at higher stress. Horizontal scan size is 20 μm.
In order to test this method it was applied to PET coated with carbon coating [35,36]. The sample was elongated up to 60% at 90 °C,
and the irreversible deformation of the carbon coating was equal to
Fig. 7 shows TEM images of Formvar films with aluminum
coatings deposited by thermal sputtering. The aluminum layer is
continuous even at coating thickness of 1.8 nm (Fig. 7a). The mosaic
structure of the coating is observed. An increase in coating thickness
up to 4 nm does not change the coating structure (Fig. 7b). Further increase of coating thickness (to 16 nm) leads to formation of crystals in
aluminum (dark spots in microphotographs) with sizes of 10–30 nm
(Figs. 7c).
The structure of the aluminum coating is similar to that of SiO2
coating deposited on PET substrate by ion sputtering deposition.
SiO2 coating was continuous even if its thickness was 2 nm [37]. In
contrast, at ion sputtering deposition of noble metals on the polymer
substrate coating islands were observed [38]. Different coating structure is attributed to their chemical activity. Chemically active metals
form continuous coating [39].
Fig. 8 demonstrates typical micrographs of a PET coated with
aluminum after elongation by 50% at 90°С. Both fragmentation and
regular wave-like folding of the coating are observed on micrographs.
The width of the coating fragment bands (Fig. 8a) and the length of
the surface wave (Fig. 8b) were measured. The coatings yield stress
σy was measured with Eq. (1). Fig. 9 shows the coatings yield stress
σy plotted against the thickness of the coating h. The aluminum
yield stress σy increases with the decrease in h up to ~ 150 MPa. If
h > 10–12 nm, the yield stress approaches the yield stress of aluminum in bulk (35 MPa) [40]. Thus, the decrease in the thickness of
aluminum coating causes an increase in the yield stress.
Fig. 10 shows the strength σ* of aluminum plotted against the
thickness of the coating. σ* was determined with Eq. (2). The strength
σ* depends on temperature of sample elongation. The calculated
coating strength is higher if the sample was elongated at room
temperature. This is explained by deformation strain-hardening of
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183
Fig. 7. TEM images of aluminum coatings deposited by thermal sputtering in vacuum on a Formvar film. The coating thickness is (a) 1.8 nm, (b) 4 nm, and, (c) 14 nm.
the metal, which is discussed below. σ* increases with the decrease in
the thickness of aluminum coating in both curves. The increase in
strength σ* of thin films is similar to an increase in the yield stress
(Figs. 8 and 9).
Fig. 11 shows fracture deformation of aluminum, εcoat, plotted
against coating thickness, h. The fracture strain of aluminum, εcoat,
increases with a decrease in h. If samples were tested at 20 °C, εcoat
reaches 140% which is much higher than the fracture strain of aluminum in bulk (40–50%) [40]. Thus, a decrease in thickness leads to
simultaneous increase in yield stress, strength and fracture strain of
aluminum coating. In contrast, in bulk solids an increase in strength
is usually accompanied by a decrease in fracture strain [41].
Substrate affects the fracture strain of metal coating. According to
[42], necking of metal film is suppressed if it is adhered to a polymeric
substrate. The substrate suppresses localization of yielding in the
neck and fracture of the coating as illustrated schematically by
Fig. 12. Hence, fracture strain of an isolated film is lower than that
of a coating adhered to a substrate. Fracture strain of thin copper
films is 1–2 % [43–46], while fracture strain of copper coating is
20%. In addition to suppression of coating necking, strengthening of
metal is explained by stopping of dislocations at the metal/polymer
interface [46].
Below the difference of coating strength values in samples tested
at temperatures 20 and 90 °С is discussed. Glass transition temperature Tg of PET is 75 °C, and at room temperature PET is in a glassy
state. In contrast, at temperature 90 °C PET is in a rubbery state. The
yield stress of glassy PET is much (by an order of magnitude) higher
than that in a rubbery state. In addition, at 20 °C in PET a neck is formed,
and the strain in neck is 275%. The strain in samples elongated at 90 °C
is only 50%.
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Fig. 10. Strength of aluminum plotted against the coating thickness. Samples were elongated at 20 (curve 1) and (2) 90 °C. Strength of block aluminum is 60–70 MPa [40].
PET coated with 9 nm-thick aluminum was elongated at 90 °C.
Strength and fracture strain of the aluminum coating were measured
independently using Eqs. (2) and (3). Bulk solids usually fracture by
break-up of a sample into two parts. In contrast, multiple fragmentation of the coatings is observed. This allows to determine strength and
fracture strain of the coating. Fig. 13 shows the coating strength plotted against its fracture strain. For comparison, the curve 2 shows the
stress–strain curve of bulk aluminum [40]. The strength of thin aluminum coating is much higher than that of bulk aluminum. The strength
of aluminum at strain of 40% increases to ~ 800 MPa, while the stress
of bulk aluminum remains practically unchanged.
Fig. 13 shows essential strain-hardening of the aluminum coating. It
may explain strength difference between the coatings elongated at
Fig. 8. SEM micrographs of PET coated with aluminum film (10 nm in thickness)
elongated at 90 °C by 50%. The arrow shows the elongation axis. The sample was tilted
by 45°.
Fig. 11. Fracture deformation of aluminum plotted against the coating thickness.
Samples were elongated by 50% at 20 (1) and 90 °C (2).
Fig. 9. Yield stress of aluminum plotted against the coating thickness. Samples were
elongated on 50% at 90 °C.
Fig. 12. Schematic of localization of yielding in an isolated metallic film (top image)
and in a coating adhered to a polymer substrate (bottom) [43].
A.L. Volynskii et al. / Thin Solid Films 536 (2013) 179–186
Fig. 13. (1) The strength of 9-nm-thick aluminum coating plotted against plastic deformation at 90 °C; (2) stress–strain curve of bulk aluminum at room temperature [31].
temperatures 20 and 90 °C (Fig. 9). The substrate at temperature 90 °C
is elongated by 50% while at 20 °C it is elongated by 275%. At 20 °C,
the strain of the polymer substrate is higher. The strain-hardening
causes higher plastic strain and higher stress of the aluminum coating.
In order to explain an increase of the strength and yield stress
with a decrease in the coatings thickness, SEM data are considered
Fig. 14. Electron diffraction images of aluminum. The coating thickness is 4 (a) and
16 nm (b).
185
(Fig. 7). The structure of an aluminum coating depends on its thickness. In relatively thick (10 nm or more) coatings, crystals are observed (Fig. 7c and d). The size of crystals is equal to 10–30 nm. In
thinner coatings mosaic structure is observed (Fig. 6a).
Electron diffraction (Fig. 14) reveals different structure of thin and
thick aluminum coatings. Crystalline structure is observed only if the
thickness of the coating is > 5 nm. In thinner coatings amorphous
halo is observed. If the coating thickness is less than 5–10 nm, the
metal is amorphous.
Methods of making of amorphous metal alloy films are described
in [47]. One of these methods consists in high-speed ion beam
sputtering of metals [48,49]. By this process nanosize coatings may
be obtained. If the thickness of the metal layer does not exceed tens
of nanometers, the layer is amorphous, and crystals are not formed.
Intermediate case between completely amorphous and crystalline
states of metals corresponds to amorphous–nanocrystalline composites. They consist of nanocrystal inclusions in amorphous metal
matrix. Such structures were studied in the mid-1980s [50,51]. They
were obtained by rapid quenching of alloys from the melt [50]
followed by annealing and partial crystallization of the amorphous
alloys [52]. Glezer et al. [50] proposed a structure model of metals
(Fig. 15) and pointed to the appearance of amorphous-crystalline
materials with very high strength. These materials are very hard
(hardness is ≈21 GPa) which is more than two-fold higher than
that (≈9 GPa) of the amorphous alloy. Tensile strength of the
composite at room temperature is very high (6–6.5 GPa). In the
same time, these materials are ductile. The theoretical model shown
in Fig. 15 explains unusually high hardness, yield stress, strength
and plasticity of this material [48]. High yield stress is explained by
low size of crystals and presence of disclinations.
According to transmission electron microscopy (Fig. 7), the structure of thin aluminum coatings is similar to that of amorphousnanocrystalline composites (Fig. 15). Thus, high strength, yield stress
and fracture strain of thin aluminum coating may be explained by the
size effect. An increase in strength and yield stress, as well as in fracture strain is related to a decrease in the thickness of the coating and
to change of metal structure.
Fig. 15. Structural model of amorphous-crystalline transition at fast cooling of metal
melt: (1) crystals with variable lattice parameter, (2) the region of smooth transition
from crystalline to amorphous state, and (3) thin amorphous interlayers [50].
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5. Conclusions
1. Experimental methods of measurement of the tensile strength,
yield stress and fracture strain of coatings were developed.
2. Tensile strength, yield stress and fracture strain of thin aluminum
coating increase with a decrease in its thicknesses. Decrease in
the coating yield stress and the strength is explained by decreasing
size of crystals and strain-hardening of aluminum.
3. Aluminum is amorphous if the coating thickness is lower than
5–10 nm.
Acknowledgments
The authors express their sincere thanks to EA Puklina and SS
Abramchuk for their kind assistance in conducting the experiments.
This work was supported by the Russian Foundation for Basic
Research, Project #10-03-90028Bel-a, grant of State Support of Leading
Scientific Schools HIII 4371.2010.3, and State contract #02.740.11.01143.
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