- Email: [email protected]
Effect of structural relaxation on plastic flow in a Ni–Nb metallic glassy film
Available online at www.sciencedirect.com
Acta Materialia 60 (2012) 3667–3676 www.elsevier.com/locate/actamat
Effect of structural relaxation on plastic flow in a Ni–Nb metallic glassy film Y. Ma a,b, Q.P. Cao a,b,⇑, S.X. Qu a,b,c, X.D. Wang a,b, J.Z. Jiang a,b,⇑ b
a International Center for New Structured Materials, Zhejiang University, Hangzhou 310027, People’s Republic of China Laboratory of New Structured Materials, Department of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, People’s Republic of China c Institute of Applied Mechanics, Zhejiang University, Hangzhou 310027, People’s Republic of China
Received 31 December 2011; accepted 6 March 2012 Available online 11 April 2012
Abstract The effect of structural relaxation on the deformation behavior of Ni60Nb40 glassy films was systematically investigated. The film and Ti substrate were co-bent and morphology evolution of the film during bending was monitored. With a decrease in thickness there was a mode change from highly localized to non-localized deformation, with the critical thickness dependent on annealing temperature. During co-bending only a proportion of the stored elastic energy in the metallic glassy film can transfer to the shear band due to the presence of the Ti substrate, which increases with annealing temperature due to the decrease in interfacial adhesion between the film and the substrate. Thus the decrease in critical thickness and the “thickness window” for deformation mode change caused by annealing arises from the enhanced hardness/Young’s modulus and the reduced interfacial adhesion. Ó 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Metallic glasses; Mechanical properties; Annealing; Shear band; Non-localized deformation
1. Introductions Metallic glasses (MGs), possessing liquid-like atomic structures without long-range atomic periodicity, exhibit attractive mechanical properties, such as high strength approaching the theoretical limit and a large elastic limit (2%), compared with their crystalline counterparts [1–3]. However, they suffer from a strong tendency for plastic strain localization in a narrow region called the shear band, and exhibit macroscopically catastrophic failure under tension at room temperature [4,5]. The plastic flow behavior of metallic glasses can be modulated by the temperature and strain rate [6–8], and such a highly localized to ⇑ Corresponding authors at: International Center for New Structured Materials, Zhejiang University, Hangzhou 310027, People’s Republic of China. Tel.: +86 571 8795 2107; fax: +86 571 8795 1528 (Q.P. Cao). E-mail addresses: [email protected] (Q.P. Cao), [email protected] (J.Z. Jiang).
non-localized deformation mode change was also achievable by decreasing the specimen size under compression and tension [9–23]. There are two remaining puzzles that need to be resolved. (1) What is the mechanism for deformation mode transition and critical sample size for such transition? It was proposed that shear band formation may require a critical strained volume, below which the groups of shear transformation zones (STZs) cannot develop into shear bands [9,11]. Such a critical length scale is of the order of 100 nm [12,24], but is still under debate [14–18]. More investigations on these topics are urgently needed. (2) Is there an intermediate deformation mode between inhomogeneous and homogeneous deformation? In order to answer this question Kim et al. [23] recently prepared nanolaminates with alternating layers of a Cu50Zr50 metallic glass and nanocrystalline Cu of different thicknesses and performed uniaxial tensile testing. It was found that the “thickness window” for the deformation mode to change from highly localized to non-localized is
1359-6454/$36.00 Ó 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.actamat.2012.03.014
3668
Y. Ma et al. / Acta Materialia 60 (2012) 3667–3676
relatively narrow, only about 46.8 nm, and an intermediate deformation mode may not exist in the metallic glass. These reported results were important for understanding the mechanical behavior of metallic glasses. However, a question still remains, i.e. does such a “thickness window” for highly localized to non-localized deformation mode change exist in monolithic metallic glasses? In this work we report the results for the “thickness window” for deformation mode change in magnetron-sputtered monolithic Ni–Nb metallic glass films. The effects of structural relaxation annealing on the deformation mode and the “thickness window” were investigated. 2. Experimental procedures Thin films with the nominal composition Ni60Nb40 (at.%) with thicknesses ranging from 50 to 1040 nm were deposited on a cleaned Si wafer and polished Ti plate separately in a direct current magnetron sputtering system (DCMS) (JZCK-400). The composition of the films studied was examined by field-emission scanning electron microscopy (FE-SEM) (Hitachi S-4800) with electron dispersive X-ray spectroscopy (EDS) and found to be on average Ni59.7Nb40.3, which is close to the nominal composition. The thickness and amorphous nature of deposited films were characterized by synchrotron radiation X-ray reflectivity (SR-XRR) and diffraction (SR-XRD) in beamline BL14B1 of the Shanghai Synchrotron Radiation Facility (SSRF) at a wavelength of 0.12398 nm, field-emission high resolution transmission electron microscopy (FEHRTEM) (JEOL 2100F) and differential scanning calorimetry (DSC) (Netzsch 404C). The hardness and elastic modulus of the thin film were measured by nanoindentation (Agilent Nano Indenter G200). Bars 2 mm wide, 1 mm thick, and 15 mm long were cut from strips deposited on a Ti substrate and bent using a mandrel with a radius of 15 mm along the width of the bar using a home-made apparatus [25]. SEM (Hitachi TM-1000, Hitachi S-4800) was used to monitor the morphology evolution of the film after bending under different conditions. 3. Results and discussion 3.1. Structural characterization and morphology evolution of as-sputtered films during bending Ni60Nb40 films of different thicknesses were prepared as shown in Fig. 1a. The thickness change versus deposition time is almost linear with a deposition rate of 17.6 ± 0.5 nm min1. The thickness of the thinnest film was measured as 53 ± 1 nm by SR-XRR [26] (Fig. 1b), consistent with the value of 50 ± 2 nm obtained from SEM observation (Fig. 1a). No sharp Bragg diffraction peaks for crystalline phases were detected by SR-XRD (see, for example, the XRD pattern for the as-deposited 50 nm film in Fig. 2a). The macroscopic morphology of bent film is shown in Fig. 2b. No lattice fringes and only a homogeneous maze
Fig. 1. (a) The film thickness as a function of the deposition time. (Inset) Typical SEM images of the fracture surface of thin films, from which the film thickness can be directly measured. (b) XRR data and its best fit for the thinnest film. From the fitting data we can accurately determine the film thickness as 53 ± 1 nm.
contrast were found by HRTEM (Fig. 2c), while the inset shows a selected area electron diffraction (SAED) pattern without crystalline spots and rings, further confirming the monolithic amorphous phase. Surface morphology evolution of the film during bending was monitored by SEM. Here the SEM observations focused on a location 0.2 mm from the tensile edge (the black circle marked in Fig. 2b), where the specimen experienced a total tensile strain of about 5%. Because the bent bars are composed of a Ti substrate and a sputtered amorphous thin film morphology evolution of the titanium substrate alone was examined during bending at room temperature at a strain rate of 2 103 s1, as shown in Fig. 3. No obvious slip bands on the surface are observed and only some random traces exist, probably due to polishing. Fig. 4a–d shows the morphology evolution of the as-sputtered thin films of different thicknesses bent at
Y. Ma et al. / Acta Materialia 60 (2012) 3667–3676
3669
Fig. 2. (a) Typical XRD pattern for an as-sputtered sample with a thickness of 50 nm detected by synchrotron radiation X-ray diffraction (k = 0.12398 nm) and for the 250 nm thick samples before and after annealing detected by common X-ray diffraction (k = 0.1542 nm). Here the 2h of the 50 nm thick sample XRD pattern has been converted for the sake of comparison. (b) The macroscopic morphology of a bent sample. (c and d) HRTEM images of the as-sputtered sample with a thickness of 50 nm and the as-annealed sample annealed at 623 K for 1 h.
Fig. 3. Morphology of Ti substrate bent to 5% tensile strain at room temperature at a strain rate of 2 103 s1.
room temperature and a strain rate of 2 103 s1. Nine samples were tested for each thickness. At thicknesses greater than 284 nm shear bands perpendicular to the
direction of applied stress are clearly observed, and some cracks even appear (Fig. 4a). With a decrease in thickness the propensity for shear banding gradually decreases and the proportion of non-localized deformation increases. In other words, in nine testing samples for each thickness the number of samples deforming in a non-localized mode increases with a decrease in thickness, as representatively shown in Fig. 4b and c for the 268 and 200 nm thick samples, respectively. On further reducing the thickness shear bands were completely absent in bent thin films with a thickness of less than 184 nm, as shown in Fig. 4d, indicating non-localized deformation. Shear bands cannot be detected for thin films with a thickness of less than 184 nm even by SEM at high magnifications of up to 100,000, as shown in Fig. 5. The number ratios of samples deforming in a nonlocalized mode in nine as-sputtered samples for each thickness are depicted in Fig. 6 as a function of thickness. The number ratio gradually increases with the decreasing thickness, with the critical thickness for the deformation mode change being about 184–284 nm, which is slightly
3670
Y. Ma et al. / Acta Materialia 60 (2012) 3667–3676
Fig. 4. Surface morphology of the (a) 284 nm, (b) 268 nm, (c) 200 nm and (d) 184 nm thick as-sputtered samples deformed at a strain rate of 2 103 s1 to a tensile strain of 5% at room temperature.
Fig. 5. A series of SEM images of the as-sputtered sample with a thickness of 184 nm at different magnifications showing the authenticity of nonlocalized deformation. There are no shear bands even at magnifications up to (a) 10,000 and (b) 100,000.
Fig. 6. The number ratio of samples deforming in non-localized mode in nine tested samples for each thickness as a function of thickness and annealing condition.
larger than the approximately 100–150 nm achieved in uniaxial tension tests of nanolaminates with a Cu–Zr amorphous layer and a Cu nanocrystalline layer [23]. We speculate that such a difference in the critical thickness
for non-localized deformation may arise from: (1) the different alloy system; (2) the influence of two boundaries between the amorphous and crystalline layers in the nanolaminate, which will depress the shear banding instability
Y. Ma et al. / Acta Materialia 60 (2012) 3667–3676
[21,22] of the metallic glass, leading to an increased critical thickness; (3) the change in loading conditions from uniaxial tensile loading (without confinement and decreasing critical thickness) to film–substrate co-bending (with constraint and probably increasing critical thickness). For instance, in the microcompression experiments the Pd–Si MG pillar had a critical diameter of 400 nm for nonlocalized deformation [11], while shear banding is still present in a Zr–Ti–Cu–Ni–Be MG pillar with a diameter as small as 150 nm [16]. Competition between these aspects may determine the critical thickness for deformation mode change in metallic glasses. 3.2. Annealing-induced microstructure, mechanical properties and deformation mode change Annealing at temperatures well below the glass transition temperature can result in obvious structural relaxation, reduce the open volume content and consequently affect the mechanical response of metallic glasses. Thus we carried out annealing treatments of magnetron-sputtered monolithic Ni–Nb metallic glass films at 523 and 623 K for 1 h (below the glass transition temperature of 891 K for a Ni–Nb metallic glass ribbon [27]). The microstructures of two annealed samples (523 and 623 K for 1 h) were observed by HRTEM. Fig. 2d shows a bright field image of the sample annealed at 623 K for 1 h, and the inset is the corresponding SAED pattern. No obvious lattice fringes and diffraction spots from crystalline phases were detected, confirming the amorphous nature of the annealed samples. The annealed samples were also bent
3671
and the morphology evolution observed. Fig. 7a–d shows the morphology evolution of annealed samples with different thicknesses bent at room temperature and 2 103 s1. For the sample annealed at 523 K for 1 h the critical thickness for complete shear banding is about 132 nm, smaller than 284 nm for the as-sputtered sample, and the morphology of the sample of this thickness is shown in Fig. 7a. With a decrease in thickness the propensity for shear banding gradually reduces and the proportion of non-localized deformation increases, as shown in Fig. 7b and c for the thicknesses 117 and 100 nm, respectively. On further reducing the thickness non-localized deformation becomes the dominant deformation mode at thicknesses of less than 84 nm, as shown in Fig. 7d, with a surface morphology very similar to that observed in Fig. 3. On increasing the annealing temperature from 523 to 623 K the minimum thickness for completely localized deformation was further reduced to 100 nm, and the maximum thickness for nonlocalized deformation fell to 50 nm, as shown in Fig. 7e and f, respectively, and also in Fig. 6. The hardness and Young’s modulus of the as-sputtered and as-annealed samples with a thickness of 1000 nm were measured by nanoindentation using the continuous stiffness measurement (CSM) technique. The CSM technique imposes a harmonic force on top of the nominally increasing indentation load on the indenter, and allows continuous measurement of contact stiffness as a function of indentation depth. Using the CSM technique unloading data is unnecessary for the analysis as the contact stiffness is obtained during loading indentation experiments. The relation between contact stiffness and indentation depth
Fig. 7. Surface morphology of the (a) 132 nm, (b) 117 nm, (c) 100 nm and (d) 84 nm thick as-annealed samples annealed at 523 K for 1 h deformed at 2 103 s1 to a tensile strain of 5% at room temperature. On annealing at 623 K for 1 h (e) shear bands appear in the 100 nm thick sample, while (f) non-localized deformation occurs in the 50 nm thick sample.
3672
Y. Ma et al. / Acta Materialia 60 (2012) 3667–3676
pffiffiffi can be expressed as S ¼ b kEr h [28], where S is the contact stiffness, b = 1.167 and k ¼ 24:56 for a Berkovich indenter tip, h is the indentation depth, Er is the reduced Young’s modulus, which can be derived as 1=E ¼ 1 v2i =Ei þ r 1 v2s =Es , where Ei and Es are defined in terms of the Young’s modulus of the indenter and sample, and vi and vs are the Poisson ratios for the indenter and sample, respectively. Ei = 1141 GPa and Poisson ratio vi = 0.07 are used for diamond indenter tips, and ms = 0.37 for the sample [29]. The maximum indentation depth was set at 190 nm. Typical indentation load–depth (P–h) curves for the as-sputtered and as-annealed samples are shown in Fig. 8a. It is clear that with an increase in annealing temperature a large indentation load is required to reach the same indentation depth, indicating that the hardness or Young’s modulus is enhanced by annealing treatment. Fig. 8b depicts the contact stiffness as a function of indentation depth and sample condition. The curves are almost linear except for some variation at greater indentation depths. From the slopes of these curves the Young’s modulus of the as-sputtered and as-annealed samples can be obtained with ms = 0.37, as shown in Fig. 8c. The hardness as a function of indentation depth is shown in Fig. 8d, deduced from the P–h curve and the already obtained
Young’s modulus of the sample. It is evident that the hardness and elastic modulus remain almost constant in the indentation depth range 70–190 nm. Normally, in order to measure “film only” properties a commonly used rule of thumb is to limit the indentation depth to less than 10% of the film thickness, and then the obtained hardness can be considered to be accurate without a contribution from the substrate. Therefore, the hardness at an indentation depth of 100 nm was selected as the hardness of the sample as a rough estimation. The hardness was 9.8, 10.2 and 10.6 GPa for the as-sputtered and 523 and 623 K annealed samples, respectively. The average Young’s modulus is about 149.8, 152.9 and 156.9 GPa for the assputtered and 523 and 623 K annealed samples, respectively. It was reported that during the early stage of film growth the coalescence of isolated islands and the removal of voids between the islands, driven by a reduction in surface free energy, can lead to tensile stress in thin films [30,31]. Recently Zhang et al. [32] and Raghavan et al. [33] controllably induced compressive residual stresses by shot peening, and found that the hardness of the shot peened layer is significantly increased compared with the undeformed matrix. Chen et al. [34] revealed that the tensile stress introduced by elastic bending can decrease the
Fig. 8. (a) Typical indentation load–depth (P–h) curves for the as-sputtered and as-annealed samples. (b) Contact stiffness as a function of indentation depth and sample condition. (c) Young’s modulus of the as-sputtered and as-annealed samples as a function of indentation depth with ms = 0.37. (d) The hardness as a function of indentation depth deduced from the P–h curve and the already obtained Young’s modulus values.
Y. Ma et al. / Acta Materialia 60 (2012) 3667–3676
measured hardness. Therefore, in this work annealing effectively increased the hardness and Young’s modulus, resulting from, most likely, not only a reduction in the open volume, but also the release of residual tensile stress by the annealing treatment. 3.3. Possible origins of deformation mode change with film thickness and annealing conditions In order to understand the origin of deformation mode change with thickness and annealing conditions it is helpful to consider the conditions required for shear band propagation. By analogy with Griffith’s crack propagation criterion, a shear band can only propagate if the elastic strain energy relief is larger than the surface energy increase due to the formation of a shear band [11]. Previous works [11,12] on the effect of specimen size on compressive and tensile deformation in metallic glasses indicated that the critical stress needed for shear band formation was associated with specimen height in the direction of applied stress. From the aspect ratio (height/diameter) the relevance of critical stress to the initiation of shear banding and the diameter of the specimen is thus established. For the SEM observed region in the bent sample the strategy proposed by Chen et al. [18] was adopted, by considering a deforming unit volume inside the thin film V = h3 that accommodates an individual shear band, where h is the film thickness. For the bending of a rectangular beam the elastic strain energy density Ue = r2/6E, where r is the maximum stress before shear banding and E is the Young’s modulus. Assuming that elastic energy is completely transferred to the shear band traversing a thin film in the thickness direction at an angel of 45°, pffiffiffi the energy per unit area of shear band C ¼ h 2r2 =12E. E of the as-sputtered sample measured by nanoindentation is about 149.8 GPa, consistent with the results of compression experiments on bulk specimen [35]. Recently Chen et al. [18] revealed that the value of C for Zr- and Cu-based metallic glasses was about 0.56 J m2, not as high as the 10 J m2 reported by Dubach et al. [17], which explains why the expected transition in deformation mode from localized to non-localized flow was not observed there [17], since the actual critical sample size is much lower than the expected value. Here C = 0.56 rather than 10 J m2 is used in this work. Assuming the ideal strength of metallic glass rth pE/30 [11,12], there exists a ffiffiffi critical thickness scale h ¼ 6 2EC=r2th below which the strength of metallic glass is equal to rth. When the specimen thickness is below this critical thickness materials are insensitive to structural flaws [36], and shear bands cannot be formed and the plastic strain becomes homogeneously distributed. Using the already obtained values for Young’s modulus from nanoindentation experiments, the critical thicknesses h in the as-sputtered and 523 and 623 K annealed samples are about 28.5, 28.0 and 27.3 nm, which is much lower than the values 184, 84 and 50 nm observed in the experiment.
3673
The assumption that the elastic energy is completely transferred to the shear band traversing a thin film in the thickness direction at an angle of 45° is often proposed when analyzing a primary shear band during deformation and fracture, but only when there is no other way to release the elastic strain energy. In the case of the present work one side of the sample is attached to a Ti substrate, and such confinement can lead to the suppression of film fracture. We speculate that only a proportion of stored elastic energy in the film will be transfered to the shear band, and such a ratio Ps may not be sensitively related to the sample thickness, since the sample thickness is much smaller than pthe ffiffiffi substrate thickness (1 mm). The equation P s h ¼ 6 2EC=r2th can be obtained considering only that part of the elastic energy transferred to the shear band. Using the data mentioned above and the experimentally observed critical thickness for non-localized deformation in the as-sputtered sample 184 nm, Ps is calculated as about 15.5%. Assuming that Ps did not change with the sample conditions (annealing treatments used here), the values of h for the 523 and 623 K annealed samples were about 180 and 176 nm, respectively, using the already obtained values for the Young’s modulus for the annealed samples obtained from the nanoindentation experiments. These two values are much larger than those observed experimentally, 84 and 50 nm. If the experimentally observed values 84 andpffiffi50 ffi nm were substituted in the equation P s h ¼ 6 2EC=r2th the values of Ps for the 523 and 623 K annealed samples are estimated to be about 33.3% and 54.5%, respectively. The results may suggest that the value of Ps could increase with annealing treatment. Therefore, the variation in interfacial adhesion between a thin film and substrate with annealing treatment needs to be confirmed. Nanoscratch testing is a common method to qualitatively measure the interfacial adhesion properties of nanoscale thin film/substrate interfaces. A loaded cube corner indenter tip was drawn through the surface under an increasing normal load, with a maximum of 10 mN, with a lateral profiling and scratch velocity of 10 lm s1. The indenter tip moved 100 lm laterally under a profiling load of 20 lN, then the normal load was linearly increased to the maximum and meanwhile the surface was scratched for a length of 500 lm. The abrupt change in the lateral load–displacement curve when interfacial delamination occurs was used to determine the critical load for failure. This critical load can be used as a qualitative measure of film/substrate adhesion. When there is no applied lateral load available abrupt fluctuations in the displacement in the normal direction caused by interfacial delamination/ failure can also be used to determine the critical load for film failure [37]. The larger the specific displacement in the normal direction, the larger the normal and lateral loads. Three scratch tests were performed for the assputtered and annealed samples. Fig. 9a shows typical normal displacement–scratch displacement curves for the as-sputtered and 523 and 623 K annealed samples with a
3674
Y. Ma et al. / Acta Materialia 60 (2012) 3667–3676
Fig. 9. (a) Typical normal displacement–scratch displacement curves for the 150 nm thick as-sputtered and 523 and 623 K annealed samples. (c and d) Corresponding SEM images of the scratch tracks in the as-sputtered and 523 and 623 K annealed samples.
thickness of 150 nm. It is clear that for the as-sputtered sample no abrupt fluctuation was observed, indicating that no interfacial delamination/failure takes place for the maximum load adopted. With the increase in annealing temperature an abrupt change in the curve occurs for the lower displacement, suggesting decreasing interfacial adhesion with increasing annealing temperature. It should be noted that for the as-sputtered sample the maximum displacement into the surface in the normal direction was much more than the thickness of the sample, but still no interfacial delamination was detected, indicating excellent interfacial adhesion. Displacement in the normal direction before the occurrence of interfacial delamination/failure was about 400 nm even for the 623 K annealed sample, still more than twice the sample thickness, indicating the success of film deformation observed experimentally during film/substrate co-bending. The corresponding SEM images of the scratch tracks in the as-sputtered and 523 and 623 K annealed samples are shown in Fig. 9b–d. No buckling was detected in the as-sputtered sample, while some obvious fracture/buckling was observed in the 523 and 623 K annealed samples. The origin of reduced interfacial adhesion after annealing treatment may be due to the
enhanced mismatch in Young’s modulus between the film and substrate. The Young’s modulus of Ti is about 110 GPa (http://en.wikipedia.org/wiki/Young%27s_modulus), which is much lower than those of the as-sputtered and as-annealed films shown in Fig. 8. With the increase in the annealing temperature the difference in modulus between the film and substrate further increases, resulting in a reduction in interfacial adhesion and enhancement of Ps. The release of residual tensile stress in thin films by annealing treatment might also modify the interfacial adhesion. The “thickness window” for the deformation mode change in metallic glass was also found to be affected by annealing treatment, as shown in Fig. 6. The “thickness window” is about 48 and 50 nm for the samples annealed at 523 and 623 K, respectively, which is evidently lower than the value of 100 nm for the as-sputtered sample. Jang and Greer [12] have proposed that the upper bound of stress for non-localized deformation is the ideal strength of a metallic glass, while the lower bound is the yield strength. Therefore, the minimum thickness for completely shear banding pffiffiffican be calculated from the equation P s h ¼ 6 2EC=r2th if the yield strength ry is known and
Y. Ma et al. / Acta Materialia 60 (2012) 3667–3676
substitutes for rth. As the hardness of a metallic glass is normally three times the yield strength [38] the values of ry are estimated to be about 3.3, 3.4 and 3.5 GPa for the as-sputtered and 523 and 623 K annealed samples, respectively. The calculated minimum thickness for completely shear banding is about 422 nm for the as-sputtered specimen with Ps = 15.5%, which is slightly larger than the value of 284 nm observed experimentally. For the annealed samples the calculated minimum thickness for completely shear banding decreases from 189 to 112 nm, which is very close to the values of 132 and 100 nm observed experimentally as the annealing temperature is increased from 523 K with Ps = 33.3% to 623 K with Ps = 54.5%. The calculated “thickness window” does indeed decrease with increasing annealing temperature, which is consistent with the experimental results, although the exact calculated and experimental values for the “thickness window” are slightly different. 4. Conclusions In summary, a comprehensive investigation into the effect of structural relaxation on the deformation mode and “thickness window” for deformation mode change in magnetron sputtered monolithic Ni–Nb metallic glass films was performed. The results allowed the following conclusions to be drawn. (1) For the as-sputtered sample the minimum thickness for complete shear banding is about 284 nm. With a decrease in thickness the propensity for shear banding gradually decreases and the proportion of nonlocalized deformation increases. The critical thickness for completely non-localized deformation is about 184 nm. The “thickness window” for a change from highly localized to non-localized deformation mode is about 100 nm, which is slightly larger than the value of about 46.8 nm achieved in uniaxial tension tests of nanolaminates with a Cu–Zr amorphous layer and a Cu nanocrystalline layer [23]. (2) Structural relaxation at 523 and 623 K for 1 h does not induce nanocrystallization, but enhances the hardness and Young’s modulus. With an increase in annealing temperature the critical thickness for complete shear banding decreases from 132 nm for 523 K annealing to 100 nm for 623 K annealing, while the critical thickness for completely non-localized deformation reduces from 84 nm for 523 K annealing to 50 nm for 623 K annealing. (3) Assuming that the elastic energy is completely transferred to a shear band traversing the thin film the enhanced Young’s modulus caused by structural relaxation will result in the calculated critical thickness for completely non-localized deformation being much lower than that observed experimentally. Thus only a proportion of the stored elastic energy is transfered to the shear band, due to the presence of the Ti substrate, and the ratio Ps is calculated to be about
3675
15.5%, 33.3% and 54.5% for the as-sputtered and 523 and 623 K annealed samples. The value of Ps increases with increasing annealing temperature due to a decrease in interfacial adhesion between the film and the substrate. (4) The “thickness window” for deformation mode change in metallic glasses is found to be affected by annealing treatment. The “thickness window” is about 48 and 50 nm for the samples annealed at 523 and 623 K, respectively, which is much lower than value of 100 nm for the as-sputtered sample. The decrease in “thickness window” in the annealed films probably arises from the larger decrease in critical thickness for complete shear banding caused by the increased hardness (yield strength) compared with the reduction in critical thickness for completely nonlocalized deformation caused by the enhanced Young’s modulus, and the reduced interfacial adhesion. Acknowledgements Financial support from the National Natural Science Foundation of China (Grants Nos. 50701038, 50920105101, 51050110136, 51071141, 10979002 and 10904127), the National Key Basic Research Program of China (2012CB825700), National High Technology Research and Development Program of China (2012AA041206), Zhejiang Provincial Natural Science Foundation of China (Y4110192), the Zhejiang University–Helmholtz Cooperation Fund, the Fundamental Research Funds for the Central Universities, and the Department of Science and Technology of Zhejiang Province are gratefully acknowledged. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19]
Greer AL. Science 1995;267:1947. Inoue A. Acta Mater 2000;48:279. Wang WH, Dong C, Shek CH. Mater Sci Eng R 2004;44:45. Chen MW. Annu Rev Mater Res 2008;38:445. Schuh CA, Hufnagel TC, Ramamurty U. Acta Mater 2007;55:4067. Kawamura YT, Nakamura T, Inoue A. Scripta Mater 1998;39:301. Schuh CA, Lund AC, Nieh TG. Acta Mater 2004;52:5879. Spaepen F. Acta Metall 1977;25:407. Li QK, Li M. Intermetallics 2006;14:1005. Guo H, Yan PF, Wang YB, Tan J, Zhang ZF, Sui ML, et al. Nat Mater 2007;6:735. Volkert CA, Donohue A, Spaepen F. J Appl Phys 2008;103:083539. Jang DC, Greer JR. Nat Mater 2010;9:215. Luo JH, Wu FF, Huang JY, Wang JQ, Mao SX. Phys Rev Lett 2010;104:215503. Schuster BE, Wei Q, Hufnagel TC, Ramesh KT. Acta Mater 2008;56:5091. Shan ZW, Li J, Cheng YQ, Minor AM, Syed Asif SA, Warren OL, et al. Phys Rev B 2008;77:155419. Wu XL, Guo YZ, Wei Q, Wang WH. Acta Mater 2009;57:3562. Dubach A, Raghavan R, Lo¨ffler JF, Michler J, Ramamurty U. Scripta Mater 2009;60:567. Chen CQ, Pei YT, De Hosson JTM. Acta Mater 2010;58:189. Bharathula A, Lee SW, Wright WJ, Flores KM. Acta Mater 2010;58:5789.
3676
Y. Ma et al. / Acta Materialia 60 (2012) 3667–3676
[20] Yavari A, Georgarakis K, Botta W, Inoue A, Vaughan G. Phys Rev B 2010;82:172202. [21] Donohue A, Spaepen F, Hoagland RG, Misra A. Appl Phys Lett 2007;91:241905. [22] Wang YM, Li J, Hamza AV, Barbee TW. Proc Natl Acad Sci USA 2007;104:11155. [23] Kim JY, Jang DC, Greer JR. Adv Funct Mater 2011;21:4550. [24] Shimizu F, Ogata S, Li J. Mater Trans 2007;48:2923. [25] Wang X, Cao QP, Chen YM, Hono K, Zhong C, Jiang QK, et al. Acta Mater 2011;59:1037. [26] Yang P, Lu D, Ramana Murthy B, Moser HO. Surf Coat Technol 2005;198:133. [27] Chen LY, Fu ZD, Zeng W, Zhang GQ, Zeng YW, Xu GL, et al. J Alloys Compd 2007;443:105. [28] Zeng K. In: Tseng TY, Nalwa HS, editors. Handbook of nanoceramics and their based nanodevices. American Scientific Publishers; 2009.
[29] [30] [31] [32] [33] [34] [35] [36] [37] [38]
Wang WH. J Appl Phys 2006;99:093506. Thornton JA, Hoffman DW. Thin Solid Films 1989;171:5. Nix WD, Clemens BM. J Mater Res 1999;14:3467. Zhang Y, Wang WH, Greer AL. Nat Mater 2006;5:857. Raghavan R, Ayer R, Jin HW, Marzinsky CN, Ramamurty U. Scripta Mater 2008;59:167. Chen LY, Ge Q, Qu S, Jiang JZ. Scripta Mater 2008;59:1210. Zhu ZW, Zhang HF, Pan DG, Sun WS, Hu ZQ. Adv Eng Mater 2006;8:953. Gao HJ, Ji BH, Jager IL, Arzt E, Fratzl P. Proc Natl Acad Sci USA 2003;100:5597. Qi J, Lai KH, Bello I, Lee CS, Lee ST, Luo JB, et al. J Vac Sci Technol A 2000;19:130. Schuh CA, Nieh TG. J Mater Res 2004;19:46.
Acta Materialia 60 (2012) 3667–3676 www.elsevier.com/locate/actamat
Effect of structural relaxation on plastic flow in a Ni–Nb metallic glassy film Y. Ma a,b, Q.P. Cao a,b,⇑, S.X. Qu a,b,c, X.D. Wang a,b, J.Z. Jiang a,b,⇑ b
a International Center for New Structured Materials, Zhejiang University, Hangzhou 310027, People’s Republic of China Laboratory of New Structured Materials, Department of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, People’s Republic of China c Institute of Applied Mechanics, Zhejiang University, Hangzhou 310027, People’s Republic of China
Received 31 December 2011; accepted 6 March 2012 Available online 11 April 2012
Abstract The effect of structural relaxation on the deformation behavior of Ni60Nb40 glassy films was systematically investigated. The film and Ti substrate were co-bent and morphology evolution of the film during bending was monitored. With a decrease in thickness there was a mode change from highly localized to non-localized deformation, with the critical thickness dependent on annealing temperature. During co-bending only a proportion of the stored elastic energy in the metallic glassy film can transfer to the shear band due to the presence of the Ti substrate, which increases with annealing temperature due to the decrease in interfacial adhesion between the film and the substrate. Thus the decrease in critical thickness and the “thickness window” for deformation mode change caused by annealing arises from the enhanced hardness/Young’s modulus and the reduced interfacial adhesion. Ó 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Metallic glasses; Mechanical properties; Annealing; Shear band; Non-localized deformation
1. Introductions Metallic glasses (MGs), possessing liquid-like atomic structures without long-range atomic periodicity, exhibit attractive mechanical properties, such as high strength approaching the theoretical limit and a large elastic limit (2%), compared with their crystalline counterparts [1–3]. However, they suffer from a strong tendency for plastic strain localization in a narrow region called the shear band, and exhibit macroscopically catastrophic failure under tension at room temperature [4,5]. The plastic flow behavior of metallic glasses can be modulated by the temperature and strain rate [6–8], and such a highly localized to ⇑ Corresponding authors at: International Center for New Structured Materials, Zhejiang University, Hangzhou 310027, People’s Republic of China. Tel.: +86 571 8795 2107; fax: +86 571 8795 1528 (Q.P. Cao). E-mail addresses: [email protected] (Q.P. Cao), [email protected] (J.Z. Jiang).
non-localized deformation mode change was also achievable by decreasing the specimen size under compression and tension [9–23]. There are two remaining puzzles that need to be resolved. (1) What is the mechanism for deformation mode transition and critical sample size for such transition? It was proposed that shear band formation may require a critical strained volume, below which the groups of shear transformation zones (STZs) cannot develop into shear bands [9,11]. Such a critical length scale is of the order of 100 nm [12,24], but is still under debate [14–18]. More investigations on these topics are urgently needed. (2) Is there an intermediate deformation mode between inhomogeneous and homogeneous deformation? In order to answer this question Kim et al. [23] recently prepared nanolaminates with alternating layers of a Cu50Zr50 metallic glass and nanocrystalline Cu of different thicknesses and performed uniaxial tensile testing. It was found that the “thickness window” for the deformation mode to change from highly localized to non-localized is
1359-6454/$36.00 Ó 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.actamat.2012.03.014
3668
Y. Ma et al. / Acta Materialia 60 (2012) 3667–3676
relatively narrow, only about 46.8 nm, and an intermediate deformation mode may not exist in the metallic glass. These reported results were important for understanding the mechanical behavior of metallic glasses. However, a question still remains, i.e. does such a “thickness window” for highly localized to non-localized deformation mode change exist in monolithic metallic glasses? In this work we report the results for the “thickness window” for deformation mode change in magnetron-sputtered monolithic Ni–Nb metallic glass films. The effects of structural relaxation annealing on the deformation mode and the “thickness window” were investigated. 2. Experimental procedures Thin films with the nominal composition Ni60Nb40 (at.%) with thicknesses ranging from 50 to 1040 nm were deposited on a cleaned Si wafer and polished Ti plate separately in a direct current magnetron sputtering system (DCMS) (JZCK-400). The composition of the films studied was examined by field-emission scanning electron microscopy (FE-SEM) (Hitachi S-4800) with electron dispersive X-ray spectroscopy (EDS) and found to be on average Ni59.7Nb40.3, which is close to the nominal composition. The thickness and amorphous nature of deposited films were characterized by synchrotron radiation X-ray reflectivity (SR-XRR) and diffraction (SR-XRD) in beamline BL14B1 of the Shanghai Synchrotron Radiation Facility (SSRF) at a wavelength of 0.12398 nm, field-emission high resolution transmission electron microscopy (FEHRTEM) (JEOL 2100F) and differential scanning calorimetry (DSC) (Netzsch 404C). The hardness and elastic modulus of the thin film were measured by nanoindentation (Agilent Nano Indenter G200). Bars 2 mm wide, 1 mm thick, and 15 mm long were cut from strips deposited on a Ti substrate and bent using a mandrel with a radius of 15 mm along the width of the bar using a home-made apparatus [25]. SEM (Hitachi TM-1000, Hitachi S-4800) was used to monitor the morphology evolution of the film after bending under different conditions. 3. Results and discussion 3.1. Structural characterization and morphology evolution of as-sputtered films during bending Ni60Nb40 films of different thicknesses were prepared as shown in Fig. 1a. The thickness change versus deposition time is almost linear with a deposition rate of 17.6 ± 0.5 nm min1. The thickness of the thinnest film was measured as 53 ± 1 nm by SR-XRR [26] (Fig. 1b), consistent with the value of 50 ± 2 nm obtained from SEM observation (Fig. 1a). No sharp Bragg diffraction peaks for crystalline phases were detected by SR-XRD (see, for example, the XRD pattern for the as-deposited 50 nm film in Fig. 2a). The macroscopic morphology of bent film is shown in Fig. 2b. No lattice fringes and only a homogeneous maze
Fig. 1. (a) The film thickness as a function of the deposition time. (Inset) Typical SEM images of the fracture surface of thin films, from which the film thickness can be directly measured. (b) XRR data and its best fit for the thinnest film. From the fitting data we can accurately determine the film thickness as 53 ± 1 nm.
contrast were found by HRTEM (Fig. 2c), while the inset shows a selected area electron diffraction (SAED) pattern without crystalline spots and rings, further confirming the monolithic amorphous phase. Surface morphology evolution of the film during bending was monitored by SEM. Here the SEM observations focused on a location 0.2 mm from the tensile edge (the black circle marked in Fig. 2b), where the specimen experienced a total tensile strain of about 5%. Because the bent bars are composed of a Ti substrate and a sputtered amorphous thin film morphology evolution of the titanium substrate alone was examined during bending at room temperature at a strain rate of 2 103 s1, as shown in Fig. 3. No obvious slip bands on the surface are observed and only some random traces exist, probably due to polishing. Fig. 4a–d shows the morphology evolution of the as-sputtered thin films of different thicknesses bent at
Y. Ma et al. / Acta Materialia 60 (2012) 3667–3676
3669
Fig. 2. (a) Typical XRD pattern for an as-sputtered sample with a thickness of 50 nm detected by synchrotron radiation X-ray diffraction (k = 0.12398 nm) and for the 250 nm thick samples before and after annealing detected by common X-ray diffraction (k = 0.1542 nm). Here the 2h of the 50 nm thick sample XRD pattern has been converted for the sake of comparison. (b) The macroscopic morphology of a bent sample. (c and d) HRTEM images of the as-sputtered sample with a thickness of 50 nm and the as-annealed sample annealed at 623 K for 1 h.
Fig. 3. Morphology of Ti substrate bent to 5% tensile strain at room temperature at a strain rate of 2 103 s1.
room temperature and a strain rate of 2 103 s1. Nine samples were tested for each thickness. At thicknesses greater than 284 nm shear bands perpendicular to the
direction of applied stress are clearly observed, and some cracks even appear (Fig. 4a). With a decrease in thickness the propensity for shear banding gradually decreases and the proportion of non-localized deformation increases. In other words, in nine testing samples for each thickness the number of samples deforming in a non-localized mode increases with a decrease in thickness, as representatively shown in Fig. 4b and c for the 268 and 200 nm thick samples, respectively. On further reducing the thickness shear bands were completely absent in bent thin films with a thickness of less than 184 nm, as shown in Fig. 4d, indicating non-localized deformation. Shear bands cannot be detected for thin films with a thickness of less than 184 nm even by SEM at high magnifications of up to 100,000, as shown in Fig. 5. The number ratios of samples deforming in a nonlocalized mode in nine as-sputtered samples for each thickness are depicted in Fig. 6 as a function of thickness. The number ratio gradually increases with the decreasing thickness, with the critical thickness for the deformation mode change being about 184–284 nm, which is slightly
3670
Y. Ma et al. / Acta Materialia 60 (2012) 3667–3676
Fig. 4. Surface morphology of the (a) 284 nm, (b) 268 nm, (c) 200 nm and (d) 184 nm thick as-sputtered samples deformed at a strain rate of 2 103 s1 to a tensile strain of 5% at room temperature.
Fig. 5. A series of SEM images of the as-sputtered sample with a thickness of 184 nm at different magnifications showing the authenticity of nonlocalized deformation. There are no shear bands even at magnifications up to (a) 10,000 and (b) 100,000.
Fig. 6. The number ratio of samples deforming in non-localized mode in nine tested samples for each thickness as a function of thickness and annealing condition.
larger than the approximately 100–150 nm achieved in uniaxial tension tests of nanolaminates with a Cu–Zr amorphous layer and a Cu nanocrystalline layer [23]. We speculate that such a difference in the critical thickness
for non-localized deformation may arise from: (1) the different alloy system; (2) the influence of two boundaries between the amorphous and crystalline layers in the nanolaminate, which will depress the shear banding instability
Y. Ma et al. / Acta Materialia 60 (2012) 3667–3676
[21,22] of the metallic glass, leading to an increased critical thickness; (3) the change in loading conditions from uniaxial tensile loading (without confinement and decreasing critical thickness) to film–substrate co-bending (with constraint and probably increasing critical thickness). For instance, in the microcompression experiments the Pd–Si MG pillar had a critical diameter of 400 nm for nonlocalized deformation [11], while shear banding is still present in a Zr–Ti–Cu–Ni–Be MG pillar with a diameter as small as 150 nm [16]. Competition between these aspects may determine the critical thickness for deformation mode change in metallic glasses. 3.2. Annealing-induced microstructure, mechanical properties and deformation mode change Annealing at temperatures well below the glass transition temperature can result in obvious structural relaxation, reduce the open volume content and consequently affect the mechanical response of metallic glasses. Thus we carried out annealing treatments of magnetron-sputtered monolithic Ni–Nb metallic glass films at 523 and 623 K for 1 h (below the glass transition temperature of 891 K for a Ni–Nb metallic glass ribbon [27]). The microstructures of two annealed samples (523 and 623 K for 1 h) were observed by HRTEM. Fig. 2d shows a bright field image of the sample annealed at 623 K for 1 h, and the inset is the corresponding SAED pattern. No obvious lattice fringes and diffraction spots from crystalline phases were detected, confirming the amorphous nature of the annealed samples. The annealed samples were also bent
3671
and the morphology evolution observed. Fig. 7a–d shows the morphology evolution of annealed samples with different thicknesses bent at room temperature and 2 103 s1. For the sample annealed at 523 K for 1 h the critical thickness for complete shear banding is about 132 nm, smaller than 284 nm for the as-sputtered sample, and the morphology of the sample of this thickness is shown in Fig. 7a. With a decrease in thickness the propensity for shear banding gradually reduces and the proportion of non-localized deformation increases, as shown in Fig. 7b and c for the thicknesses 117 and 100 nm, respectively. On further reducing the thickness non-localized deformation becomes the dominant deformation mode at thicknesses of less than 84 nm, as shown in Fig. 7d, with a surface morphology very similar to that observed in Fig. 3. On increasing the annealing temperature from 523 to 623 K the minimum thickness for completely localized deformation was further reduced to 100 nm, and the maximum thickness for nonlocalized deformation fell to 50 nm, as shown in Fig. 7e and f, respectively, and also in Fig. 6. The hardness and Young’s modulus of the as-sputtered and as-annealed samples with a thickness of 1000 nm were measured by nanoindentation using the continuous stiffness measurement (CSM) technique. The CSM technique imposes a harmonic force on top of the nominally increasing indentation load on the indenter, and allows continuous measurement of contact stiffness as a function of indentation depth. Using the CSM technique unloading data is unnecessary for the analysis as the contact stiffness is obtained during loading indentation experiments. The relation between contact stiffness and indentation depth
Fig. 7. Surface morphology of the (a) 132 nm, (b) 117 nm, (c) 100 nm and (d) 84 nm thick as-annealed samples annealed at 523 K for 1 h deformed at 2 103 s1 to a tensile strain of 5% at room temperature. On annealing at 623 K for 1 h (e) shear bands appear in the 100 nm thick sample, while (f) non-localized deformation occurs in the 50 nm thick sample.
3672
Y. Ma et al. / Acta Materialia 60 (2012) 3667–3676
pffiffiffi can be expressed as S ¼ b kEr h [28], where S is the contact stiffness, b = 1.167 and k ¼ 24:56 for a Berkovich indenter tip, h is the indentation depth, Er is the reduced Young’s modulus, which can be derived as 1=E ¼ 1 v2i =Ei þ r 1 v2s =Es , where Ei and Es are defined in terms of the Young’s modulus of the indenter and sample, and vi and vs are the Poisson ratios for the indenter and sample, respectively. Ei = 1141 GPa and Poisson ratio vi = 0.07 are used for diamond indenter tips, and ms = 0.37 for the sample [29]. The maximum indentation depth was set at 190 nm. Typical indentation load–depth (P–h) curves for the as-sputtered and as-annealed samples are shown in Fig. 8a. It is clear that with an increase in annealing temperature a large indentation load is required to reach the same indentation depth, indicating that the hardness or Young’s modulus is enhanced by annealing treatment. Fig. 8b depicts the contact stiffness as a function of indentation depth and sample condition. The curves are almost linear except for some variation at greater indentation depths. From the slopes of these curves the Young’s modulus of the as-sputtered and as-annealed samples can be obtained with ms = 0.37, as shown in Fig. 8c. The hardness as a function of indentation depth is shown in Fig. 8d, deduced from the P–h curve and the already obtained
Young’s modulus of the sample. It is evident that the hardness and elastic modulus remain almost constant in the indentation depth range 70–190 nm. Normally, in order to measure “film only” properties a commonly used rule of thumb is to limit the indentation depth to less than 10% of the film thickness, and then the obtained hardness can be considered to be accurate without a contribution from the substrate. Therefore, the hardness at an indentation depth of 100 nm was selected as the hardness of the sample as a rough estimation. The hardness was 9.8, 10.2 and 10.6 GPa for the as-sputtered and 523 and 623 K annealed samples, respectively. The average Young’s modulus is about 149.8, 152.9 and 156.9 GPa for the assputtered and 523 and 623 K annealed samples, respectively. It was reported that during the early stage of film growth the coalescence of isolated islands and the removal of voids between the islands, driven by a reduction in surface free energy, can lead to tensile stress in thin films [30,31]. Recently Zhang et al. [32] and Raghavan et al. [33] controllably induced compressive residual stresses by shot peening, and found that the hardness of the shot peened layer is significantly increased compared with the undeformed matrix. Chen et al. [34] revealed that the tensile stress introduced by elastic bending can decrease the
Fig. 8. (a) Typical indentation load–depth (P–h) curves for the as-sputtered and as-annealed samples. (b) Contact stiffness as a function of indentation depth and sample condition. (c) Young’s modulus of the as-sputtered and as-annealed samples as a function of indentation depth with ms = 0.37. (d) The hardness as a function of indentation depth deduced from the P–h curve and the already obtained Young’s modulus values.
Y. Ma et al. / Acta Materialia 60 (2012) 3667–3676
measured hardness. Therefore, in this work annealing effectively increased the hardness and Young’s modulus, resulting from, most likely, not only a reduction in the open volume, but also the release of residual tensile stress by the annealing treatment. 3.3. Possible origins of deformation mode change with film thickness and annealing conditions In order to understand the origin of deformation mode change with thickness and annealing conditions it is helpful to consider the conditions required for shear band propagation. By analogy with Griffith’s crack propagation criterion, a shear band can only propagate if the elastic strain energy relief is larger than the surface energy increase due to the formation of a shear band [11]. Previous works [11,12] on the effect of specimen size on compressive and tensile deformation in metallic glasses indicated that the critical stress needed for shear band formation was associated with specimen height in the direction of applied stress. From the aspect ratio (height/diameter) the relevance of critical stress to the initiation of shear banding and the diameter of the specimen is thus established. For the SEM observed region in the bent sample the strategy proposed by Chen et al. [18] was adopted, by considering a deforming unit volume inside the thin film V = h3 that accommodates an individual shear band, where h is the film thickness. For the bending of a rectangular beam the elastic strain energy density Ue = r2/6E, where r is the maximum stress before shear banding and E is the Young’s modulus. Assuming that elastic energy is completely transferred to the shear band traversing a thin film in the thickness direction at an angel of 45°, pffiffiffi the energy per unit area of shear band C ¼ h 2r2 =12E. E of the as-sputtered sample measured by nanoindentation is about 149.8 GPa, consistent with the results of compression experiments on bulk specimen [35]. Recently Chen et al. [18] revealed that the value of C for Zr- and Cu-based metallic glasses was about 0.56 J m2, not as high as the 10 J m2 reported by Dubach et al. [17], which explains why the expected transition in deformation mode from localized to non-localized flow was not observed there [17], since the actual critical sample size is much lower than the expected value. Here C = 0.56 rather than 10 J m2 is used in this work. Assuming the ideal strength of metallic glass rth pE/30 [11,12], there exists a ffiffiffi critical thickness scale h ¼ 6 2EC=r2th below which the strength of metallic glass is equal to rth. When the specimen thickness is below this critical thickness materials are insensitive to structural flaws [36], and shear bands cannot be formed and the plastic strain becomes homogeneously distributed. Using the already obtained values for Young’s modulus from nanoindentation experiments, the critical thicknesses h in the as-sputtered and 523 and 623 K annealed samples are about 28.5, 28.0 and 27.3 nm, which is much lower than the values 184, 84 and 50 nm observed in the experiment.
3673
The assumption that the elastic energy is completely transferred to the shear band traversing a thin film in the thickness direction at an angle of 45° is often proposed when analyzing a primary shear band during deformation and fracture, but only when there is no other way to release the elastic strain energy. In the case of the present work one side of the sample is attached to a Ti substrate, and such confinement can lead to the suppression of film fracture. We speculate that only a proportion of stored elastic energy in the film will be transfered to the shear band, and such a ratio Ps may not be sensitively related to the sample thickness, since the sample thickness is much smaller than pthe ffiffiffi substrate thickness (1 mm). The equation P s h ¼ 6 2EC=r2th can be obtained considering only that part of the elastic energy transferred to the shear band. Using the data mentioned above and the experimentally observed critical thickness for non-localized deformation in the as-sputtered sample 184 nm, Ps is calculated as about 15.5%. Assuming that Ps did not change with the sample conditions (annealing treatments used here), the values of h for the 523 and 623 K annealed samples were about 180 and 176 nm, respectively, using the already obtained values for the Young’s modulus for the annealed samples obtained from the nanoindentation experiments. These two values are much larger than those observed experimentally, 84 and 50 nm. If the experimentally observed values 84 andpffiffi50 ffi nm were substituted in the equation P s h ¼ 6 2EC=r2th the values of Ps for the 523 and 623 K annealed samples are estimated to be about 33.3% and 54.5%, respectively. The results may suggest that the value of Ps could increase with annealing treatment. Therefore, the variation in interfacial adhesion between a thin film and substrate with annealing treatment needs to be confirmed. Nanoscratch testing is a common method to qualitatively measure the interfacial adhesion properties of nanoscale thin film/substrate interfaces. A loaded cube corner indenter tip was drawn through the surface under an increasing normal load, with a maximum of 10 mN, with a lateral profiling and scratch velocity of 10 lm s1. The indenter tip moved 100 lm laterally under a profiling load of 20 lN, then the normal load was linearly increased to the maximum and meanwhile the surface was scratched for a length of 500 lm. The abrupt change in the lateral load–displacement curve when interfacial delamination occurs was used to determine the critical load for failure. This critical load can be used as a qualitative measure of film/substrate adhesion. When there is no applied lateral load available abrupt fluctuations in the displacement in the normal direction caused by interfacial delamination/ failure can also be used to determine the critical load for film failure [37]. The larger the specific displacement in the normal direction, the larger the normal and lateral loads. Three scratch tests were performed for the assputtered and annealed samples. Fig. 9a shows typical normal displacement–scratch displacement curves for the as-sputtered and 523 and 623 K annealed samples with a
3674
Y. Ma et al. / Acta Materialia 60 (2012) 3667–3676
Fig. 9. (a) Typical normal displacement–scratch displacement curves for the 150 nm thick as-sputtered and 523 and 623 K annealed samples. (c and d) Corresponding SEM images of the scratch tracks in the as-sputtered and 523 and 623 K annealed samples.
thickness of 150 nm. It is clear that for the as-sputtered sample no abrupt fluctuation was observed, indicating that no interfacial delamination/failure takes place for the maximum load adopted. With the increase in annealing temperature an abrupt change in the curve occurs for the lower displacement, suggesting decreasing interfacial adhesion with increasing annealing temperature. It should be noted that for the as-sputtered sample the maximum displacement into the surface in the normal direction was much more than the thickness of the sample, but still no interfacial delamination was detected, indicating excellent interfacial adhesion. Displacement in the normal direction before the occurrence of interfacial delamination/failure was about 400 nm even for the 623 K annealed sample, still more than twice the sample thickness, indicating the success of film deformation observed experimentally during film/substrate co-bending. The corresponding SEM images of the scratch tracks in the as-sputtered and 523 and 623 K annealed samples are shown in Fig. 9b–d. No buckling was detected in the as-sputtered sample, while some obvious fracture/buckling was observed in the 523 and 623 K annealed samples. The origin of reduced interfacial adhesion after annealing treatment may be due to the
enhanced mismatch in Young’s modulus between the film and substrate. The Young’s modulus of Ti is about 110 GPa (http://en.wikipedia.org/wiki/Young%27s_modulus), which is much lower than those of the as-sputtered and as-annealed films shown in Fig. 8. With the increase in the annealing temperature the difference in modulus between the film and substrate further increases, resulting in a reduction in interfacial adhesion and enhancement of Ps. The release of residual tensile stress in thin films by annealing treatment might also modify the interfacial adhesion. The “thickness window” for the deformation mode change in metallic glass was also found to be affected by annealing treatment, as shown in Fig. 6. The “thickness window” is about 48 and 50 nm for the samples annealed at 523 and 623 K, respectively, which is evidently lower than the value of 100 nm for the as-sputtered sample. Jang and Greer [12] have proposed that the upper bound of stress for non-localized deformation is the ideal strength of a metallic glass, while the lower bound is the yield strength. Therefore, the minimum thickness for completely shear banding pffiffiffican be calculated from the equation P s h ¼ 6 2EC=r2th if the yield strength ry is known and
Y. Ma et al. / Acta Materialia 60 (2012) 3667–3676
substitutes for rth. As the hardness of a metallic glass is normally three times the yield strength [38] the values of ry are estimated to be about 3.3, 3.4 and 3.5 GPa for the as-sputtered and 523 and 623 K annealed samples, respectively. The calculated minimum thickness for completely shear banding is about 422 nm for the as-sputtered specimen with Ps = 15.5%, which is slightly larger than the value of 284 nm observed experimentally. For the annealed samples the calculated minimum thickness for completely shear banding decreases from 189 to 112 nm, which is very close to the values of 132 and 100 nm observed experimentally as the annealing temperature is increased from 523 K with Ps = 33.3% to 623 K with Ps = 54.5%. The calculated “thickness window” does indeed decrease with increasing annealing temperature, which is consistent with the experimental results, although the exact calculated and experimental values for the “thickness window” are slightly different. 4. Conclusions In summary, a comprehensive investigation into the effect of structural relaxation on the deformation mode and “thickness window” for deformation mode change in magnetron sputtered monolithic Ni–Nb metallic glass films was performed. The results allowed the following conclusions to be drawn. (1) For the as-sputtered sample the minimum thickness for complete shear banding is about 284 nm. With a decrease in thickness the propensity for shear banding gradually decreases and the proportion of nonlocalized deformation increases. The critical thickness for completely non-localized deformation is about 184 nm. The “thickness window” for a change from highly localized to non-localized deformation mode is about 100 nm, which is slightly larger than the value of about 46.8 nm achieved in uniaxial tension tests of nanolaminates with a Cu–Zr amorphous layer and a Cu nanocrystalline layer [23]. (2) Structural relaxation at 523 and 623 K for 1 h does not induce nanocrystallization, but enhances the hardness and Young’s modulus. With an increase in annealing temperature the critical thickness for complete shear banding decreases from 132 nm for 523 K annealing to 100 nm for 623 K annealing, while the critical thickness for completely non-localized deformation reduces from 84 nm for 523 K annealing to 50 nm for 623 K annealing. (3) Assuming that the elastic energy is completely transferred to a shear band traversing the thin film the enhanced Young’s modulus caused by structural relaxation will result in the calculated critical thickness for completely non-localized deformation being much lower than that observed experimentally. Thus only a proportion of the stored elastic energy is transfered to the shear band, due to the presence of the Ti substrate, and the ratio Ps is calculated to be about
3675
15.5%, 33.3% and 54.5% for the as-sputtered and 523 and 623 K annealed samples. The value of Ps increases with increasing annealing temperature due to a decrease in interfacial adhesion between the film and the substrate. (4) The “thickness window” for deformation mode change in metallic glasses is found to be affected by annealing treatment. The “thickness window” is about 48 and 50 nm for the samples annealed at 523 and 623 K, respectively, which is much lower than value of 100 nm for the as-sputtered sample. The decrease in “thickness window” in the annealed films probably arises from the larger decrease in critical thickness for complete shear banding caused by the increased hardness (yield strength) compared with the reduction in critical thickness for completely nonlocalized deformation caused by the enhanced Young’s modulus, and the reduced interfacial adhesion. Acknowledgements Financial support from the National Natural Science Foundation of China (Grants Nos. 50701038, 50920105101, 51050110136, 51071141, 10979002 and 10904127), the National Key Basic Research Program of China (2012CB825700), National High Technology Research and Development Program of China (2012AA041206), Zhejiang Provincial Natural Science Foundation of China (Y4110192), the Zhejiang University–Helmholtz Cooperation Fund, the Fundamental Research Funds for the Central Universities, and the Department of Science and Technology of Zhejiang Province are gratefully acknowledged. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19]
Greer AL. Science 1995;267:1947. Inoue A. Acta Mater 2000;48:279. Wang WH, Dong C, Shek CH. Mater Sci Eng R 2004;44:45. Chen MW. Annu Rev Mater Res 2008;38:445. Schuh CA, Hufnagel TC, Ramamurty U. Acta Mater 2007;55:4067. Kawamura YT, Nakamura T, Inoue A. Scripta Mater 1998;39:301. Schuh CA, Lund AC, Nieh TG. Acta Mater 2004;52:5879. Spaepen F. Acta Metall 1977;25:407. Li QK, Li M. Intermetallics 2006;14:1005. Guo H, Yan PF, Wang YB, Tan J, Zhang ZF, Sui ML, et al. Nat Mater 2007;6:735. Volkert CA, Donohue A, Spaepen F. J Appl Phys 2008;103:083539. Jang DC, Greer JR. Nat Mater 2010;9:215. Luo JH, Wu FF, Huang JY, Wang JQ, Mao SX. Phys Rev Lett 2010;104:215503. Schuster BE, Wei Q, Hufnagel TC, Ramesh KT. Acta Mater 2008;56:5091. Shan ZW, Li J, Cheng YQ, Minor AM, Syed Asif SA, Warren OL, et al. Phys Rev B 2008;77:155419. Wu XL, Guo YZ, Wei Q, Wang WH. Acta Mater 2009;57:3562. Dubach A, Raghavan R, Lo¨ffler JF, Michler J, Ramamurty U. Scripta Mater 2009;60:567. Chen CQ, Pei YT, De Hosson JTM. Acta Mater 2010;58:189. Bharathula A, Lee SW, Wright WJ, Flores KM. Acta Mater 2010;58:5789.
3676
Y. Ma et al. / Acta Materialia 60 (2012) 3667–3676
[20] Yavari A, Georgarakis K, Botta W, Inoue A, Vaughan G. Phys Rev B 2010;82:172202. [21] Donohue A, Spaepen F, Hoagland RG, Misra A. Appl Phys Lett 2007;91:241905. [22] Wang YM, Li J, Hamza AV, Barbee TW. Proc Natl Acad Sci USA 2007;104:11155. [23] Kim JY, Jang DC, Greer JR. Adv Funct Mater 2011;21:4550. [24] Shimizu F, Ogata S, Li J. Mater Trans 2007;48:2923. [25] Wang X, Cao QP, Chen YM, Hono K, Zhong C, Jiang QK, et al. Acta Mater 2011;59:1037. [26] Yang P, Lu D, Ramana Murthy B, Moser HO. Surf Coat Technol 2005;198:133. [27] Chen LY, Fu ZD, Zeng W, Zhang GQ, Zeng YW, Xu GL, et al. J Alloys Compd 2007;443:105. [28] Zeng K. In: Tseng TY, Nalwa HS, editors. Handbook of nanoceramics and their based nanodevices. American Scientific Publishers; 2009.
[29] [30] [31] [32] [33] [34] [35] [36] [37] [38]
Wang WH. J Appl Phys 2006;99:093506. Thornton JA, Hoffman DW. Thin Solid Films 1989;171:5. Nix WD, Clemens BM. J Mater Res 1999;14:3467. Zhang Y, Wang WH, Greer AL. Nat Mater 2006;5:857. Raghavan R, Ayer R, Jin HW, Marzinsky CN, Ramamurty U. Scripta Mater 2008;59:167. Chen LY, Ge Q, Qu S, Jiang JZ. Scripta Mater 2008;59:1210. Zhu ZW, Zhang HF, Pan DG, Sun WS, Hu ZQ. Adv Eng Mater 2006;8:953. Gao HJ, Ji BH, Jager IL, Arzt E, Fratzl P. Proc Natl Acad Sci USA 2003;100:5597. Qi J, Lai KH, Bello I, Lee CS, Lee ST, Luo JB, et al. J Vac Sci Technol A 2000;19:130. Schuh CA, Nieh TG. J Mater Res 2004;19:46.