Comparative study on plasticity and fracture ...

Tribology International 126 (2018) 344–351

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Comparative study on plasticity and fracture behaviour of Ti/Al multilayers a,∗

b

Kunkun Fu , Leigh Sheppard , Li Chang a b

a,∗∗

a

b

, Xianghai An , Chunhui Yang , Lin Ye

T

a

Centre for Advanced Materials Technology, School of Aerospace, Mechanical and Mechatronic Engineering, The University of Sydney, NSW 2006, Australia School of Computing, Engineering and Mathematics, Western Sydney University, NSW 2751, Australia

A R T I C LE I N FO

A B S T R A C T

Keywords: Nanostructured Ti/Al multilayers Plasticity Fracture Work of adhesion

This paper presented a comparative study on plasticity and fracture behaviour of nanostructured Ti/Al multilayers with two individual layer thicknesses λ. The nanoindentation results showed that film buckling appears in the 10 nm multilayer, whereas localised shear bands are detected in the 100 nm multilayer. The nanoscratch results indicated that localised conformal cracks and spalling are the dominant fracture mode in the multilayer with λ = 10 nm, whereas adhesive failure is observed in the 100 nm multilayer when a critical normal load is reached. The work of adhesion of the 10 nm multilayer was much larger than that with λ = 10 nm, possibly due to the high hardness and few dislocation slips by the confined layer mechanism.

1. Introduction The repeating deposition of two or more materials with nanometric thickness can produce an ultra-strong nanostructured multilayer (NM). In particular, NMs made of metallic materials have been of significant interests to extensive researchers [1–9] because of their notable strength and hardness. Meanwhile, the metallic-based NMs generally show better fracture resistance than ceramic-based NMs [10–12] due to the ductile nature of most metals. These excellent properties endow metallic NMs with great applied potential as protective coatings. It has been agreed that metallic NMs exhibit length-scale strength. As the individual layer thickness of the metallic NMs is decreased from hundreds of nanometres to a few nanometres, the dominant dislocation activities alter from dislocation pile-up at interfaces to dislocation slip in confined layers [13]. When the individual layer thickness is further reduced to less than 2 nm, the dominant mechanism alters again to interface crossing by a single dislocation [14]. The variation of strengthening mechanism suggests dissimilar plastic deformation behaviour in the metallic NMs with various ranges of individual layer thickness. To unveil the deformation behaviour of metallic NMs, extensive studies [15–20] have been performed utilising the nanoindentation technique. For instance, Li et al. [15] investigated length-scale plastic instabilities in Au/Cu multilayers with different individual layer thickness using nanoindentation. In the work of Wang et al. [16], the plastic instabilities of Au/Ti multilayers revealed using a Berkovich indenter changed from shear bands to delamination when the layer thickness was increased from 25 nm to 250 nm. In our previous work [17], co-deformation behaviour was found in Al/Ti ∗

Corresponding author. Corresponding author. E-mail addresses: [email protected] (K. Fu), [email protected] (L. Chang).

∗∗

https://doi.org/10.1016/j.triboint.2018.05.033 Received 22 February 2018; Received in revised form 9 May 2018; Accepted 21 May 2018 Available online 22 May 2018 0301-679X/ © 2018 Elsevier Ltd. All rights reserved.

multilayers when the layers were compressed to less than 30 nm by a conical indenter. During nanoindentation, the stress-state in metallic NMs is mainly compressive [18]. Therefore, the indentation-induced fracture is rarely reported in ductile metallic NMs [19]. Yang et al. [19] examined the fracture behaviour of Cu/Ni/W multilayers using nanoindentation. They found that cracks initiate at brittle W layer and were arrested by Cu/W interfaces and Ni/W interfaces. Compared to nanoindentation, nanoscratch is a testing technique which produces a severe deformation in the tested samples. The nanoscratch technique has been widely used to evaluate the adhesion and fracture resistance of thin films [21–24]. To our knowledge, examination of the fracture or adhesion of metallic NMs using nanoscratch has rarely been reported [25]. Topić et al. [25] studied the effect of annealing temperature on scratch resistance in single, double and triple layers of Pt/V coated systems, and observed spallation-buckling in all test scenarios. Unfortunately, the length scale fracture behaviour of Pt/V NMs was not covered in their study [25]. In the current paper, we are the first to report the distinct fracture behaviour of metallic NMs with different individual layer thickness using nanoscratch. First, a direct current magnetron sputtering method was used to prepare Ti/Al multilayers with two individual layer thickness, i.e. 10 nm and 100 nm. Second, the plasticity behaviour of the two Ti/Al multilayers were compared using nanoindentation. Then, the fracture behaviour of the prepared Ti/Al multilayers were investigated using nanoscratch. Accordingly, the work of adhesion of the Ti/Al multilayers was determined. Finally, the fracture mechanism of the Ti/Al multilayers was discussed.


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Fig. 1. Cross-sectional view of λ = 10 nm Ti/Al multilayer by SEM and HRTEM.

2. Material preparation and experimental procedures

displacement was continuously recorded with time. The nanoscratch tests were repeated three times for each specimen. After the nanoscratch testing, the surface profiles of the two Ti/Al multilayers were examined by high-resolution SEM (Zeiss Ultra).

Ti/Al multilayers with two different individual layer thicknesses, i.e. 10 nm and 100 nm, were deposited on a monocrystalline silicon (100) substrate (University wafer, USA) using a direct current magnetron sputtering method (Orion 5, AJA International, USA) at room temperature. The Ti/Al multilayers were achieved by consecutive deposition using a titanium target (99.9%) and an aluminium target (99.9%) positioned above a rotating substrate holder. The deposition conditions and procedure were described in detail in our previous report [26]. For all the samples, the Ti layer was the last deposited to minimise the surface oxidation. The overall thickness of the Ti/Al multilayers was estimated as one micron by the deposition rate of Ti and Al. After deposition, the sample surface exhibited a brilliant metallic silver colour. Here, the cross-sectional microstructures of the Ti/ Al multilayers were detected using a scanning electron microscopy (SEM, Zeiss Auriga) and a high-resolution transmission electron microscopy (HRTEM, JEOL 3000) operated at 300 kV. The SEM and HRTEM samples were prepared using a focused ion beam (FIB) micromachining system (Zeiss Auriga). Room temperature nanoindentation experiments were carried out using a nanoindentation testing facility (Hysitron Triboindenter®) equipped with a 45° half-included diamond sphero-conical indenter (tip radius: 207 nm) by a load control method. As specified by the manufacturer, the resolution of load transducer and displacement transducer was 1 nN and 0.04 nm respectively. The peak indentation load was set as 18 mN. During nanoindentation, the loading rate and unloading rate remained the same for all the samples with a constant value of 1.8 mN/ s. The tests were repeated five times for each specimen. After nanoindentation, the cross-sectional deformation in the Ti/Al multilayers was obtained by FIB (Zeiss Auriga) and observed by SEM (Zeiss Auriga) and HRTEM (JEOL 3000). The same diamond sphero-conical indenter used for the nanoindentation tests was employed for the nanoscratch experiments. During the nanoscratch tests, the normal load of the stylus was linearly ramped from zero to the maximum within 40 s, and the normal load and depth of the stylus were measured simultaneously. Two peak normal loads were used to examine the load effect on the fracture behaviour of the Ti/Al multilayers with values of 1 mN and 4 mN respectively. For the lateral displacement in the scratch direction, a 40 μm scratching trace was applied in all scratch tests, and the lateral

3. Results and discussion 3.1. Microstructure and plasticity in Ti/Al multilayers by nanoindentation Fig. 1 shows the cross-sectional microstructure of the prepared Ti/ Al multilayer with an individual layer thickness λ of 10 nm after imaging by SEM and TEM. From the SEM image, it is seen that the Ti and Al layers are repeatedly sputtered on a silicon substrate with ultra-low roughness (< 1 nm). The sputtered Ti and Al layers deposited on the silicon substrate are relatively flat with highly uniform individual layers. By measuring the thickness of the multilayered film, it is found that the thickness ratio of Ti layer and Al layer is approximately the same, while the overall thickness is ∼1 μm, consistent with our estimated value before sputtering. Furthermore, it is worthwhile noting that in practice, achieving atomically sharp interfaces when depositing two different thin materials is difficult due to lattice mismatch [27]. However, in the inserted TEM image, it is evident that interfaces achieved between the Ti and Al layers are sharp and only minimally intermixed. The TEM image also reveals that the Ti layer has a hexagonal close-packed phase with an orientation of Ti-(0002) while the Al layer has a face-centred cubic phase with a fibre orientation of Al-(111). The TEM observation is consistent with the XRD result [26]. Contrary to the argument made by Banerjee et al. [28], phase transition between Ti and Al layers has not been observed in the present case where the individual layer thickness has been reduced to ∼10 nm after magnetron sputtering. The nanoindentation experiments were then conducted on the Ti/Al multilayered films, and the corresponding indentation load-depth relations were recorded. These are shown in Figs. 2(a) and 3(a) for λ = 10 nm and λ = 100 nm respectively. The maximum indentation depth in the multilayer with λ = 10 nm is lower than that with λ = 100 nm. By the indentation load-depth curve, the indentation hardness H can be obtained using the following equation proposed by Oliver and Pharr [29];

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Fig. 2. (a) Nanoindentation load-depth curve and (b) its cross-sectional view of deformation in Ti/Al multilayers with an individual layer thickness of 10 nm.

H=

Pmax Ac

relatively small, indicating the good repeatability of the nanoindentation tests. It should be noted that the indentation hardness determined by Eq. (1) is the composite hardness with a substrate effect, and the values obtained can only be used for parallel comparison. Two discontinuities, which are called ‘pop-ins’, are observed in the indentation load-depth curve for λ = 10 nm in Fig. 2(a) whereas no detectable

(1)

where Pmax is the peak indentation load, and Ac is the contact area. The indentation hardness of the Ti/Al multilayer with λ = 10 nm is consequently determined to be 9.29 ± 0.14 GPa, which is 87.6% higher than when λ = 100 nm. The standard deviation of the test hardness is 346


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Fig. 3. (a) Nanoindentation load-depth curve and (b) its cross-sectional view of deformation in Ti/Al multilayers with an individual layer thickness of 100 nm.

discontinuity was observed in Fig. 3(a) for λ = 100 nm. The ‘pop-in’ events in the nanoindentation load-depth curves are also repeatable in the Ti/Al multilayer with λ = 10 nm. The first ‘pop-in’ occurs when the indentation depth reaches to 388 ± 28 nm, while the second ‘pop-in’ happens at the indentation depth of 485 ± 36 nm ‘Pop-in’ events are related to a sudden energy release during nanoindentation, which is usually attributed to cracking [30] or phase transformation [31]. By

examining indentation-induced deformation in the Ti/Al multilayers, the mechanism of ‘pop-in’ formation can be identified. In Fig. 2(a), the inserted SEM image shows two ring-like wrinkles that have formed on the edge of the indent. In contrast, no wrinkles can be seen on the film surface displayed in Fig. 3(a). Fig. 2(b) shows the cross-sectional view of the indentation-induced deformation. It can be seen that the wrinkles on the film surface are actually layer buckling-induced two-fold pile347


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Fig. 4. Surface morphology in Ti/Al multilayer with an individual layer thickness of (a) 10 nm and (b) 100 nm after nanoscratch with a peak load of 1 mN.

its formation. Hence, ‘pop-in’ is not observed as shown in Fig. 3(a). Furthermore, it should be pointed out that the fracture or cracking is not detected in both Ti/Al multilayers under nanoindentation, although the maximum indentation depth exceeds half of the overall thickness of the Ti/Al multilayers. There are several studies [15,16] explaining the mechanism of distinct plastic instabilities under nanoindentation. For the Ti/Al multilayer with λ = 100 nm, the main mechanism is dislocation pile-up at interfaces. Thus, the applied shear stress concentrates at the Ti/Al interfaces during nanoindentation. Also, it has been reported [15] that the stress concentration by the dislocation pileup is directly proportional to the individual layer thickness. Therefore, for the multilayer with an elevated λ (100 nm), the stress concentration is greater than the interface strength, leading to a localised shearing. In contrast, dislocation pileup is incapable of generating high stress at the interfaces of the

ups. Two curved shear bands are found originating from the pile-up and propagate into the multilayered film. The shape of the shear bands indicates that they are formed by layer rotation. Furthermore, the inserted TEM images in Fig. 2(b) show that no obvious intermixing is observed after layer buckling. The event of layer buckling can be correlated to the ‘pop-in’ in the indentation load-depth curve as seen in Fig. 2(a). When a critical load is reached, the rotated layer buckles and the stored energy releases, resulting in a discontinuity in the indentation load-depth curves. However, in Fig. 3(b), the pile-up and shear bands in the Ti/Al multilayer with λ = 100 nm is different than those seen in the multilayer with λ = 10 nm. The pile-up is only one fold, and its height is relatively low, implying a more moderate plastic deformation in the film. Beneath the pile-up, a localised shear band is detected. The shear band is induced by the interface failure instead of layer rotation. The shear band is small with less energy released during 348


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Conformal cracks, which are defined as micro-cracks that form while the surface is attempting to conform to the scratch groove, are a typical damage feature that occurs in brittle coatings during scratching. At the end of the scratch, delaminated layers and spalling stack up after withdrawing the indenter and an abrupt change is found in the normal depth and lateral displacement curve. This is clearly evident in Fig. 5(a) for the Ti/Al multilayer with λ = 10 nm, and it conclusively shows brittle behaviour during scratching for this multilayer. In terms of the multilayer with λ = 100 nm, a clear groove is seen after a low ramping load scratching as shown in Fig. 4(b), although the normal depth is much deeper than that in the λ = 10 nm multilayer. The peak normal depth is much lower than the overall thickness of the multilayer and no exposure of the silicon substrate is found during scratching. Additionally, the SEM image shows that no cracks and debris are detected in the scratch trace, with no abrupt change in normal depth either (see Fig. 5(b)). The stacked pile-up is found at the edge of the groove, and the layers do not peel off during the ramping-load scratch process, even at the end of the scratch where the indenter was withdrawn. The smooth trace and stacked pile-up around the edge indicate that this Ti/ Al multilayer undergoes a ductile deformation process and has high scratch resistance at a low ramping load (1 mN). When applying a high ramping load (4 mN), examination of the scratch trace in Fig. 6 reveals that the failure mode in the Ti/Al multilayer with λ = 10 nm remains the same, whereas the multilayer with λ = 100 nm exhibits completely different damage behaviour. It is worthwhile noting that the maximum normal depth in Fig. 5 is still lower than the overall film thickness indicating that the stylus does not reach the silicon substrate during any scratches. For the multilayer with λ = 10 nm in Fig. 6(a), the conformal cracks, spalling and delaminated layers predominate. The normal depth initially increases with an increase of lateral displacement but then an abrupt drop is seen when the lateral displacement reaches ∼9 μm as shown in Fig. 5(a). This point is regarded as the critical point where the damage starts to occur. The normal load at this point was defined as the critical load when the damage occurs, and this definition has been widely adopted by many researchers [21–25] over the years. The loadcontrol feedback system enables the continuous increase in the normal load. When a failure occurs, a sudden energy releases, leading to an unstable state of scratching. If the detached materials are stuck underneath the indenter, the local stiffness of the system increases, resulting in a decrease in the normal depth. Moreover, the density of conformal cracks is higher than that with a low ramping load. However, when regarding the Ti/Al multilayer with λ = 100 nm, the main dominant behaviour has changed from plasticity to adhesive delamination as shown in Fig. 6(b). The SEM image reveals that the scratch trace is initially smooth with a transition to adhesive delamination at approximately halfway along the scratch trace. At this point, a discontinuity is detected in the normal depth and lateral displacement curve. Adhesive failure is usually considered to be when the layers detach from the layer beneath them by lifting or separation, and this point is regarded as a critical point to failure. At the end of the scratch groove, fish-scale-like delaminated layers are found which imply a more severe deformation process. Still, no cracks and less spalling are seen in the multilayer with λ = 100 nm, which indicates that this multilayer exhibits greater ductile than the multilayer with λ = 10 nm. Fig. 7 shows the normal load as a function of normal depth with a peak load of 4 mN. Based on a comparison of the surface profiles of the respective scratch traces of each multilayer, the critical load to failure Lc of both Ti/Al multilayers can be obtained. It is found that the Lc in Ti/Al multilayers with λ = 10 nm and λ = 100 nm has values of 721.2 ± 120.0 μN and 1856.2 ± 259.9 μN, respectively. It is evident that for these films, a high λ contributes to a lower Lc, which is the opposite tendency observed for CrN/NbN ceramic-based NMs [32]. Ceramic-based NMs exhibit brittle behaviour evidenced by the predominated cohesive failure modes. The number of CrN/NbN interfaces can effectively suppress the crack propagation. In our metallic Ti/Al NMs however, the reduction in the individual layer thickness involves a

Fig. 5. Normal depth and lateral displacement curves of Ti/Al multilayer with an individual layer thickness of (a) 10 nm and (b) 100 nm during nanoscratch.

Ti/Al multilayer with λ = 10 nm due to the confined layer slip mechanism. Thus, the shear stress is lower than the interface strength, leading to layer rotation and eventually layer buckling. In other words, the distinct plasticity behaviour in the Ti/Al multilayers depends on whether the stress concentration at Ti/Al interfaces exceeds the interface strength. 3.2. Fracture in Ti/Al multilayers by nanoscratch This section reports nanoscratch experiments that were conducted to investigate the fracture behaviour of the Ti/Al multilayers. Fig. 4 illustrates the surface morphologies of the Ti/Al multilayers with λ = 10 nm and λ = 100 nm after a low ramping load nanoscratch. In Fig. 4(a), damage can be attributed to multiple conformal crack formation and delamination. In the initial 12 μm of the scratch, the trace is relatively smooth, and the normal depth increases with less fluctuation during scratching as shown in Fig. 5(a). Afterwards conformal cracks and the associated debris start to appear. Correspondingly, large fluctuations in normal depth are observed as displayed in Fig. 5(a). 349


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Fig. 6. Surface morphology in Ti/Al multilayer with an individual layer thickness of (a) 10 nm and (b) 100 nm after nanoscratch with a peak load of 4 mN.

calculated to be 41.0 ± 1.1 J/m2 and 454.8 ± 8.9 J/m2, respectively. The work of adhesion for the λ = 100 nm Ti/Al multilayers is much higher than when λ = 10 nm. This implies that the Ti/Al multilayer with a thicker individual layer thickness exhibits a better resistance to fracture. Therefore, when designing a metallic multilayer, the optimal individual layer thickness is a compromise between fracture resistance and mechanical strength. There are two possible explanations for the distinct fracture behaviour in the Ti/Al multilayers. First, the hardness of the Ti/Al multilayer with λ = 10 nm is much greater than that with λ = 100 nm. Generally, the high strength/hardness nanostructured materials exhibit lower ductility [34]. According to the observed failure mode, the reduction in individual layer thickness results in the ductile Ti/Al multilayer changing to a multilayer with a brittle behaviour. This ductile-

ductile-to-brittle transition, with the mechanism to be discussed later. The work of adhesion W of the Ti/Al multilayers can be determined using the equation [33] as follows;

W=

32Lc2 t π 2Ec dc4

(2)

where t is the film thickness, Ec is the film modulus, and dc is the width of the scratch trace at critical load. In this case, the work of adhesion corresponds to the interface between Ti and Al layers, and not the film and substrate. Therefore, t can be approximated by the normal depth at critical load. The Rule of Mixtures is used here to determine the film modulus, which was subsequently found to be Ec = 95 GPa dc is measured from the SEM images displayed in Fig. 6. Using Eq. (2), the work of adhesion of the Ti/Al multilayers with λ = 10 nm and λ = 100 nm is 350


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Fig. 7. Normal load and depth relations in Ti/Al multilayers after nanoscratch with a peak load of 4 mN.

to-brittle transition phenomenon is observed for the first time in this investigation. Second, as already mentioned, the thin Ti and Al layers tend to rotate, bend and eventually buckle because the applied shear stress can not overcome the interface strength. When the layers are buckled, the bent Ti and Al layers are under tension at their extremities. This is where cracks are more likely to originate. In addition, discrete dislocation gliding dominates the plastic deformation process due to the limited dislocation capacity when λ = 10 nm. In contrast, the thicker layer thickness can accommodate much greater dislocation motion. Hence, there is more dislocation storage capability at a higher length scale, which promotes ductile behaviour under nanoscratch. 4. Conclusions In the present work, the plasticity and fracture behaviour of the Ti/ Al multilayers with λ = 10 nm and λ = 100 nm have been compared using nanoindentation and nanoscratch experiments respectively. The following conclusions have been drawn: (1) Film buckling occurred in the λ = 10 nm Ti/Al multilayer after nanoindentation, whereas a localised shear band was seen in the Ti/ Al multilayer with λ = 100 nm. This is determined by the competition between stress concentration induced by dislocation pileup, and interface strength. (2) Localised cracks, debris and delaminated layers were seen in Ti/Al multilayer with λ = 10 nm, whereas adhesive failure was observed in the multilayer with λ = 100 nm when a high normal load was applied. (3) Critical load to failure in Ti/Al metallic multilayer with λ = 10 nm is lower than that with λ = 100 nm. The tendency is opposite to that in ceramic multilayers. (4) The averaged work of adhesion in Ti/Al multilayers with λ = 10 nm and λ = 100 nm was calculated to be 41.0 J/m2 and 454.8 J/m2, respectively. The Ti/Al multilayer with a reduced λ has a low fracture resistance. Acknowledgements The authors greatly acknowledge the financial support from the Faculty Research Cluster Program, the University of Sydney. The scientific and technical input and support from the Australian Microscopy and Microanalysis Research Facility node at the University of Sydney are also acknowledged. 351