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Nanocrystalline electrodeposited Ni: microstructure and tensile properties
Acta Materialia 50 (2002) 3957–3970 www.actamat-journals.com
Nanocrystalline electrodeposited Ni: microstructure and tensile properties F. Dalla Torre a, H. Van Swygenhoven a,∗, M. Victoria b b
a Paul Scherrer Institute, CH-5232 Villigen, Switzerland CRPP-Fusion Technology Materials, EPFL, CH-5232 Villigen, Switzerland
Received 2 April 2002; accepted 15 May 2002
Abstract The microstructure of commercially available nanocrystalline (nc) electroplated Ni foils is studied by means of Xray diffraction and transmission electron microscopy. It is shown that the microstructure is inhomogeneous and batchdependent. Tensile properties at strain rates between 10⫺5 and 103 s⫺1 are studied and compared with the results of coarse-grained Ni. Data on strength, strain-rate sensitivity and work hardening are presented. At the highest strain rates, shear banding with local grain growth is observed in the nc structure. It is also suggested that the differences found in nc Ni for 3 and 20 mm tensile specimens are the size effects related to the inhomogeneous microstructure. 2002 Published by Elsevier Science Ltd on behalf of Acta Materialia Inc. Keywords: Nanocrystalline; Mechanical properties; Microstructure; Nickel; Shear bands
1. Introduction It is well known that in polycrystalline metals, an increase in hardness and strength can be obtained by reducing the grain size to the nanometer scale. There is, however, no clear understanding of the deformation mechanisms that operate at these grain sizes. It is expected that plasticity is, to a lesser extent, driven by dislocation mechanisms and that deformation will be accommodated in the grain boundaries. Computer simulations in full 3D nanostructures [1,2] as well as in columnar ∗ Corresponding author. Paul Scherrer Institute, Building WHGA/343, CH-5232 Villigen PSI, Switzerland. Tel.: +41-56310-21-11; fax: +41-56-310-21-99. E-mail address: [email protected] (H. Van Swygenhoven).
nanostructures [3] suggested grain boundary sliding as the main deformation process [1]. However, direct experimental observation of deformation mechanism at atomic scale is still missing. In situ deformation studies using transmission electron microscopes (TEM) [4,5] do not provide convincing results since possible relaxations in the nanosized grain boundaries due to the thinning procedure tend to obscure interpretations [6]. Another concern, and major obstacle, in the investigation of the deformation mechanism is the limited ductility reported in these nanostructures [7], which are believed to be mainly provoked by pre-existing flaws, such as impurities and porosity [8], introduced during the fabrication of the material. Therefore, a lot of attention is paid to electrodeposited (ED) and severe plastic deformed
1359-6454/02/$22.00 2002 Published by Elsevier Science Ltd on behalf of Acta Materialia Inc. PII: S 1 3 5 9 - 6 4 5 4 ( 0 2 ) 0 0 1 9 8 - 2
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nanocrystalline (nc) samples [9–12] as it is believed that these techniques can produce porosity-free samples. In recent years, efforts have been made to relate the nanostructure to the deposition parameters in order to obtain texture-free nanometer-sized grain structures with high chemical purity [13,14]. The preferential growth of crystals and the resulting texture are indeed common problems in the electrodeposition of metals and are difficult to control. Therefore, reports on mechanical properties have to be accompanied by an accurate microstructural investigation. Usually grain size and microstrain are provided by means of X-ray diffraction (XRD), but often little information is available on TEM observations [11–13]. Another major difficulty involved in performing mechanical tests on nc metals is often the limited amount of material provided by the above-mentioned techniques. Electrodeposition does not suffer from this restriction, although the sample thickness usually does not exceed 0.3 mm. As other methods, such as inert gas condensation and ball milling, do provide only a limited amount of material, miniaturised tensile machines have been developed allowing the comparison of the mechanical properties of samples synthesised by different methods [15,16]. Recent computer simulation results [1,17,35] have provided an analysis of the deformation mechanism. These simulations, though, are performed at very high strain rates. Part of the present experimental effort includes testing at high strain rates, up to 103 s⫺1, in order to extrapolate to the ultra-high strain rates reached in the simulations. This article provides a detailed description of the microstructure of a commercially available ED Ni and its room temperature mechanical behaviour under uniaxial tension. The microstructure of the ED Ni samples is characterised by means of their mean grain size, internal strains and texture deduced from XRD and TEM. Different batches are compared in order to investigate the reproducibility of the electrodeposition technique. Tensile tests are performed at strain rates ranging from 10⫺5 to 103 s⫺1 and are compared with the behaviour of coarse-grained Ni.
2. Experimental procedures Two batches (A and B) of electroplated nc Ni foils with a purity of 99.95% and a cold rolled coarse-grained Ni foil (purity 99.98%) were purchased from Goodfellow Metals. The thickness of the foils is 200–250 µm for the nc Ni and 1 mm for the coarse-grained Ni. Tensile specimen of coarsegrained Ni were cut along the rolling direction, mechanically polished to a thickness of approximately 0.5 mm and afterwards annealed at 700 °C. Chemical analysis on both the nc Ni batches was made via inductively coupled plasma optical emission spectrometry (ICP-OES). The density of the batches A and B was determined by the Archimedes immersion method using a Mettler microgram balance and diethyl-phthalate as a reference liquid. For both the batches, a density of 99.5 ± 0.5% was obtained. XRD measurements were performed with a Siemens D500 X-ray detector. Average grain size and lattice strain were calculated using the Warren– Averbach approach after Krill and Birringer [18] for the {111} / {222} and {200} / {400} family planes. Texture was measured using the relative intensities of the {111} and {200} peaks normalised over a standard powder diffraction pattern of (100% random) Ni powder. Conventional TEM investigations were made with a Jeol 2010 operated at 200 kV. TEM specimens were electropolished with a TENUPOL 5 of Struers. A solution of 3% perchloric acid, 30% ethylenglycol and 67% methanol at 15 V and ⫺10 °C was used. Grain boundaries of approximately 500 grains were delineated manually with adobe illustrator software on digitised TEM bright field images to determine the grain size distribution, which is fitted with a lognormal function. Because of the high lattice distortion of the microstructure, bright field images were preferred over the dark field images for the measurement of grain size. The high lattice distortion in the grains leads to a blurred diffraction contrast resulting in smaller grain sizes when measured by dark field images. Tensile tests were performed at room temperature with different tensile machines: (1) a strain rate controlled miniaturised tensile machine using 3 mm long dog-bone samples [16] at a strain rate
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of 2 × 10⫺4s⫺1; (2) two conventional tensile testing machines (Zwick and Schenk) using 20 mm dogbone shaped specimens at strain rates between 10⫺5 and 10⫺1 s⫺1 [19]; and (3) a Roell Amsler, type 1852, servo-hydraulic high-speed tensile machine designed for dynamic tension and compression tests with a maximal cross-head speed of 12 m s⫺1. When the elastic tensile slope is compared with the one measured in the Zwick or Schenk machine, it appears that the stiffness of this machine is 2.3 times lower. Therefore, strain values of the Roell Amsler machine were divided by this factor, so that strains can be compared. The yield strength and the plastic strain are calculated using the intercept of the stress–strain curve with the 0.2% total strain offset parallel to the elastic slope. The tensile specimens, shown in Fig. 5, have been machined by a wire Electro-DischargingMachine (EDM). The 3 mm long tensile specimens have a gauge length of 1.72 mm, thickness, approximately 0.20 mm and width, approximately 0.25 mm. The 20 mm tensile samples have a gauge cross-section of 2.5 × 0.2 mm and a gauge length of 6 mm. Local heat generated by the 0.1 mm thick wire during the spark erosion caused local grain growth to a 10 µm thick layer at the surfaces. Therefore, the specimens were additionally electropolished during 30 s using the same conditions as used for producing the TEM discs to remove the altered surface. Microhardness measurements were performed with a Vickers pyramid indenter, installed on a Zeiss Jenaphot 2000 optical microscope and were monitored by a computer. Mean hardness values are calculated from 4 to 9 µm indentation depths. Fractured surfaces of tensile specimens and indentations were examined by scanning electron microscopy using a Topcon instrument.
3. Microstructure 3.1. Chemical composition Table 1 shows the chemical analysis of the main impurities (ppm) for batches A and B. Metallic impurities in ED samples usually originate from
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the impurities in the anode material and from the chemicals used for the electrolytic bath. Cobalt and iron, for instance, are common impurities in nickel sulphate, sulphur originates from bath additives, such as natrium-saccharin (Na-C7H4NO3S) [13]. Light element impurities such as H, N and O could not be measured, but are expected to be present in the material as well. Co-deposition of hydrogen generated by the dissociation of water as well as oxidation processes in the solution is a common chemical by-product during electrodeposition. Molecules of additives, such as Na-Saccharin often used as grain-growth inhibitors may be adsorbed by the surface. Natter et al. [13] detected 110, 620 and 300 ppm hydrogen, nitrogen and oxygen, respectively, in a 17 nm grain size sample. The comparison of the amount of impurities between the batches A and B shows significant difference for some elements. The Cu content in batch B is almost three times higher. Also Al, Co and W are present in higher amounts in batch B compared to batch A. The sulphur content of batch A is, however, 15% higher. Also higher concentration of Ca and Pb are found in batch A. The differences in impurity content, indicating the difficulty in reproducing identical compositions, could play a non-negligible role in the material properties. 3.2. Texture, grain size and grain size distribution Table 2 summarises the microstructural observations of texture, grain size and strain. Compared to a standard Ni sample with a fully random crystal orientation, both the batches show preferential orientation along the {200} family planes but to different degrees. Batch A shows a 50% random orientation whereas batch B shows 30% random orientation. Data obtained from the solution side and the substrate side of both the batches give similar results. Polishing the samples results in a slightly higher observed {200} texture. TEM observations also reveal a textured microstructure. Fig. 1 shows TEM bright field images and the corresponding selected area diffraction (SAD) patterns for batch A (left) and batch B (right). The closest first ring of the diffraction pattern corresponds to the {111} peak and the second
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Table 1 Chemical impurity content of the batches A and B Element
Batch A (ppm)
Batch B (ppm)
6.5 8.75 881 155 127 19 883 ⬍
45.9 2.56 104 473 135 ⬍ 746 263
Al Ca Co Cu Fe Pb S W
Detection limit (ppm) 1 0.02 0.5 0.1 0.1 10 5 10
Table 2 Summary of the obtained microstructural measurements on grain size, strain and texture using XRD and TEM XRD {111}/ {222}
Batch A Batch B
TEM
Grain size (nm)
Microstrain (%)
Texture
Mean grain size (nm)
Peak value (nm)
Peak width
20.8 ± 3.0 21.4 ± 4.2
0.403 ± 0.015 0.393 ± 0.070
50% ± 6 30% ± 4
21.3 ± 7.8 20.6 ± 8.6
17.9 16.1
0.36 0.38
Fig. 1. TEM bright field image of batches A (left) and B (right) together with the corresponding SAD patterns (750 nm ring diameter).
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to the {200} peak of the Ni fcc structure. The difference in the relative intensities of the two rings in batches A and B demonstrates the stronger {200} texture for batch B. No additional peaks could be measured indicating that no second phase is present in the ED Ni foils. Mean grain sizes and microstrain measured with XRD reveal similar values for batches A and B. When using the {111} / {200} reflections, the corresponding values for batch A are a mean size of 20.8 ± 3.0nm and a microstrain of 0.40% whereas for batch B they are 21.4 ± 3.0nm with 0.39% strain. When the {200} / {400} reflections are used, again no notable difference between A and B is found, but here, due to the low {400} peak intensity, a large scattering between different measurements is observed (12–60 nm for the mean grain size and 0.54–0.98% for the microstrain). Grain size calculated from bright field TEM images resulted in similar mean grain sizes: 21.3 ± 7.8nm for batch A and 20.6 ± 7.8nm for batch B. Fig. 2 shows the grain size distributions for both the batches fitted with a lognormal distribution function. For batch A (B), a peak value of 17.9 nm (16.1 nm) and a peak width of 0.36 (0.38) are obtained. The ratios between the maximal and the minimal grain diameters are 1.60 for batch A and 1.63 for batch B. 3.3. Microstructural investigation by TEM TEM bright field images of both the batches (Fig. 1) show regions where grain shapes and grain
Fig. 2.
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boundaries can be clearly observed along with other regions of high density of diffraction contrast spatial variation. It is assumed that these contrasts originate mainly from microstrains, as suggested by XRD. This makes grain boundary features difficult to distinguish from other contrasts, such as bending contours and thickness fringes. Moreover, groups of grains possessing only small misorientation angles can be observed, as is shown in the higher magnification bright field image in Fig. 3a. Grains can be distinguished by their change in contrast at the grain boundaries. However, some grain boundaries show weak contrast. Tilting experiments demonstrate the presence of groups of grains with small misorientation angles. The Moire´ fringes observed in Fig. 3a can probably be ascribed to the superposition of such grains. A similar observation was made by Wang et al. [20], where the authors further suggest that these regions of small misfit are the nucleation sites of abnormal grain growth. The observation of the microstructure at high magnification revealed no clear evidence for the presence of dislocations. It should be noted that their contrasts could be convoluted with the contrast originated from microstrain in such a way that their identification is impeded. The stronger {200} texture, observed in batch B, can be associated with fibre-like structures as shown in the bright field image and the corresponding SAD pattern in Fig. 3b. The smeared-out {200} diffraction spots correspond to the dark
Grain size distributions for batches A (a) and B (b) using TEM bright field pictures fitted to a lognormal distribution.
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Fig. 3. TEM micrographs of batch B at higher magnifications: (a) region where grains can be distinguished by the change of contrast at grain boundaries. Moire´ fringes are due to the superposition of grains lying on the top of each other; (b) region with fibre-like texture. The inset shows the SAD pattern (300 nm) indicating the {200} texture.
wavy fibre-like structures in the bright field image. The fibre-like structures are inhomogeneously distributed with distances typically from 100 to 1000 nm between the groups. Fibre-type structures have also been observed in Ref. [12]. The TEM defocusing technique [21] revealed the presence of small voids with sizes between 1 and 2 nm, often situated close to or at the grain boundaries, as shown in Fig. 4 (taken from batch B). Their origin lies probably in hydrogen being incorporated during electrodeposition. The presence of voids or cavities is also confirmed by positron lifetime measurements, where the lifetimes of 160 and 350 ps are found, which relate to single vacancies and voids that are approximately 1 nm in size, respectively [22].
4. Mechanical behaviour Fig. 5a–c shows photographs of typical tensile specimens used in this study. Fig. 5b shows a 3 mm sample for batches A and B before and after the deformation at 10⫺4 s⫺1 strain rate. Examples of the 20 mm samples are shown in Fig. 5c: before testing (1), deformed at 10⫺4 s⫺1 (2), and deformed
Fig. 4. TEM image of batch B showing pores of 1–2 nm in size. The inset shows pore concentrations in the grain interior.
at 103 s⫺1 (3). In Section 4.1, the differences in mechanical behaviour between the batches are discussed. The influence of samples size of batch B is discussed in Section 4.2, and the effect of strain rate is discussed in Section 4.3.
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for the yield strength and 2.0 GPa (1.8 GPa) for the ultimate tensile strength (UTS). Both the batches showed post-necking elongation. Microhardness tests showed a similar trend as observed for the yield strength in tension—a lower hardness for batch B. The mean hardness values are 6.0 ± 0.2 and 5.5 ± 0.5GPa for batches A and B, respectively. Fig. 6 shows a Hall–Petch plot, representing our results for batches A (circles) and B (diamond) as well as the values taken from the literature [23,24]. The grain sizes are determined by XRD (full datapoints) and by the peak value of the lognormal distribution function obtained by TEM (open datapoints). 4.2. Influence of specimen size Fig. 5. Tensile specimens. (a) The size difference between big and miniaturised specimens. (b) The 3 mm specimens before testing and deformation at 10⫺4 s⫺1 for batches A and B. (c) The 20 mm specimens before testing (1) and after deformation at 10⫺4 s⫺1 (2), and at 103 s⫺1 (3). The arrow indicates the shear band, which has an orientation of about 60° with respect to the tensile axis.
Comparative tests were performed to investigate the influence of sample size in tensile testing, as
4.1. Differences in tensile properties between batches A and B Table 3 summarises the results obtained from the tensile tests, nanoindentation and microhardness measurements for both the batches. The tensile tests were performed on the 3 mm samples at a strain rate of 2 × 10⫺4s⫺1. Higher ductility and yield strength were measured for batch A. The following data are obtained for batch A (batch B): 4% (2.6%) for the plastic strain, 1.6 GPa (1.4 GPa)
Fig. 6. Hall–Petch plot representing our results among data found in literature.
Table 3 Summary of the measured mechanical properties (results of 3 mm tensile specimens) Tensile tests UTS (GPa) Batch 2.04 ± 0.10 A Batch 1.82 ± 0.09 B
Microhardness sy (GPa)
Plastic strain (%)
Hv (GPa)
1.60 ± 0.10
4.0 ± 0.2
6.0 ± 0.3
1.43 ± 0.09
2.6 ± 0.6
5.5 ± 0.1
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samples revealed the presence of necking over the whole cross-section and in most cases a homogenous rough morphology at the fractured surface indicating ductile fracture. The 20 mm samples, however, show the presence of two types of morphologies, the same rough structure and a columnar structure (see Section 3.3). It suggests that these different morphologies also represent a difference in the crystallographic orientation and/or in the shape of grains. Whether the columnar structure can be related to the fibre-type texture observed by TEM is a matter of further research. Fig. 7. Tensile curves for 3 and 20 mm specimens of batch B. Yield and UTS are lower for bigger specimens.
shown in the stress–strain plot in Fig. 7. Table 4 shows the UTS obtained for different strain rates of six tensile tests carrried out for the 3 and the 20 mm specimens. Mean values of 1849 ± 51 and 1388 ± 61 MPa were obtained for the two specimen sizes. For single crystal Ni, however, no difference was measured between the two types of samples. This means that the reason for this apparent size effect has to be found in microstructural effects, such as the presence of surface imperfections and/or the inhomogeneity of the microstructure in the gauge length. Careful examination of the EDM-cut surfaces after polishing revealed occasionally the presence of very small microcracks. Accumulation of the effect of these features is sample size dependent and could contribute to an overall lower deformation stress. On the other hand, examination of deformed 3 mm
4.3. Influence of strain rate Fig. 8 shows UTS as a function of strain rate for both nc and coarse-grained Ni. Instead of plotting the stress at a constant strain (as it is usually done for similar plots), the UTS was favoured since only few data points of stress at a certain amount of strain show overlap for both types of materials. The strain-rate sensitivity m, defined as m ⫽ (∂lns / ∂lne˙ )e, where s is the true stress, e˙ , the strain rate and e, the true strain, which is, however, calculated between 2 and 3% strain for nc and between 2 and 20% strain for coarse-grained Ni. The condition for stable deformation can be formulated [25] in terms of the strain-rate sensitivity and the work hardening rate Θ defined as Θ ⫽
Table 4 Measurements of the UTS for 3 and 20 mm specimens 3 mm tensile specimens 20 mm tensile specimens UTS Strain rate (s⫺1) (MPa)
UTS (MPa)
Strain rate (s⫺1)
7 × 10⫺5 2 × 10⫺4 2 × 10⫺4 2 × 10⫺3 1 × 10⫺2 1 × 10⫺2
1448 1397 1278 1396 1437 1374
5.5 5.5 5.5 5.5 5.5 5.5
1823 1828 1818 1921 1905 1800
× 10⫺5 × 10⫺4 × 10⫺4 × 10⫺3 × 10⫺3 × 10⫺2
Fig. 8. UTS versus the logarithm of strain rate for nc and coarse-grained Ni. The grey bar around the highest strain rate data points of coarse-grained Ni marks the uncertain data.
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1 (∂s / ∂e)e˙ . For materials with low strain-rate sensiσ tivity, plastic instability is reached when Θⱕ1. 4.3.1. Low strain rates Fig. 9a,b shows true stress–true strain curves and the work hardening rate Θ measured on 20 mm specimens of batch B using four different strain rates in the lower strain rate regime (10⫺5–10⫺2 s⫺1). Symbols are used in this figure to distinguish between curves and do not represent specific data points. As can be seen in Fig. 9, the maximal strains that can be obtained, decrease with increasing strain rate, while the UTS remains approximately constant. The UTS is, however, up to 4.5
Fig. 9. (a) True stress–true strain curve for nc Ni at strain rates in the lower strain rate regime. (b) Work hardening rate ⍜ versus true stress for the different strain rates.
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times higher (Fig. 8) in nc Ni. Plastic strain ranges from 2.2% for the lowest strain rate to 0.4% for the highest. The strain-rate sensitivity m is one order of magnitude higher for the nc (0.01–0.03) as compared to the coarse-grained Ni (0.001–0.004) for strain values between 2 and 20% total strain. The hardening rate (Fig. 9b) increases when strain rate increases from 5.5 ± 10⫺2 to 5.5 ± 10⫺5s⫺1. When compared with the coarse-grained material, a value of ⌰ ⫽ 10 at 1360 MPa is obtained from the ED Ni whereas ⌰ ⫽ 14 at 130 MPa is obtained for the coarse-grained Ni. Both represent values at 2% plastic strain. 4.3.2. High strain rates At higher strain rates the UTS increases for both the materials, the increase in the nc material is, however, more pronounced, up to 70% (2500 MPa) when strain rate is increased to 103 s⫺1. The strainrate sensitivity at these high stress levels is 0.47. No value for m could be calculated for the coarsegrained Ni, since the data at high strain rates (above 4 × 102s⫺1) have to be taken with caution, as discussed subsequently. Similar to the Kolsky bar technique [26], where a bullet is shot to a rod that propagates the compressive stress wave onto the specimen, the highspeed tensile machine used in this study transfers the kinetic energy of the accelerated plunger onto the specimen by an instantaneous applied force that propagates as a tensile stress wave in the specimen load train. With increasing loading velocity, the appearance of one or more unloading regions is observed during the tests (see tensile curve at 8.1 × 102s⫺1 in Fig. 10). This is interpreted as the arrival of the first reflection of the tensile wave on the specimen, which is now a compressive (unloading) stress wave. This wave superimposes on the initial tensile wave, resulting in a non-trivial wave shape. Fourier transformation analysis of the tensile data showed two main oscillations with frequencies of approximately 5.4 and 6.4 Hz and waves with these frequencies are seen to continue to propagate after the failure of the specimen. Using low pass filtering to skip the undesired higher frequencies, a continuous straight loading curve can be extracted, but the stress amplitude value is also affected. The total strain on the other
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Fig. 10. Engineering stress–strain curves at high strain rates for coarse-grained Ni.
Fig. 11. Engineering stress–strain curves at high strain rates for nc Ni.
hand, could be measured on the fractured specimens and this corresponds well with the data given by the strain gauges in the machine. In other words, the UTS at strain rates above 4 × 102s⫺1 are to be taken with caution and represent only trends. The data for nc Ni are, however, correct, since all the samples broke before the tensile wave was reflected. Fig. 10 shows a plot of the engineering stress– strain curves of coarse-grained Ni for high strain and rates ranging between 5.4 × 10⫺1 2 ⫺1 8.1 × 10 s . The maximal observed true stress value is 0.47 GPa and the maximal measured plastic strain is 17% of which 2–3% contributes to post-necking elongation. Higher strain rates result in an increase in strength. The effect on the work hardening is, however, less pronounced for strain rates between 5 × 101 and 1.3 × 102s⫺1. A more pronounced increase in ⍜ for higher strain rates can be expected in coarse-grained Ni [27]. All tested coarse-grained Ni samples show uniform deformation over the whole gauge length before the necking sets in and leads to failure. At a strain rate of 8 × 102s⫺1, oscillations are observed in the stress–strain curve, representing the reflections of the tensile wave. Fig. 11 shows the engineering stress–strain curves at high strain rates for nc Ni, measured on the 20 mm samples. Plastic strains range between 1.2 and 1.8%, but there is, however, no clear relation with the applied strain rate. At the highest strain rates, the yield point is less pronounced,
however, stress values of 2.5 GPa are supported by the nanostructure. The work hardening rate seems to increase with increasing strain rate, the values are, however, difficult to compare since plastic instability sets in very early. Values of ⌰ ⫽ 1 are reached already before 1% plastic strain. 4.4. Microstructural observation of fracture surfaces and deformed areas 4.4.1. Low strain rates The 20 mm specimens have a fractured surface perpendicular to the tensile direction and this is true for all the strain rates up to 102 s⫺1 (Fig. 5c). Minor sign of plasticity could be detected on the flat specimen surface along with the little evidence of necking. An SEM picture of some details of the fractured surface is given in Fig. 12. Within the same fractured surface, two different structures could be observed, marked by the letters R and C. The zone R represents a very rough and irregular surface structure, whereas zone C shows columnar fractured lines perpendicular to the foil plane. In this picture, the two regions are relatively small sized, but the dimensions of zones R and C can vary quite strongly, both reaching till 50% of the fractured surface. SEM examinations showed clearly that only the structure responsible for the rough fractured surface contributed to the observed plasticity. In samples, where a large percentage of the fractured surface was of type R, local reduction in gauge cross-section was observed, indicating
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Fig. 12. SEM micrographs of the surface of a fractured 20 mm nc Ni specimen tested at a low strain rate. The fractured surface shows two types of morphologies indicated by the letters R (rough) and C (columnar).
‘local’ necking. Most of the fractured surfaces of the 3 mm samples, deformed at low strain rates, showed only the morphologies of type R, as can be seen in Fig. 13. As was mentioned before, those samples showed post-necking elongation. The difference in fracture mechanism is probably related to the inhomogeneous microstructure revealed in TEM. 4.4.2. High strain rates At higher strain rates, a clear change in fracture mode is observed in the 20 mm samples. The angle of the fracture surface with the tensile direction changed from ~90 to 65–55°, corresponding to the
angle of maximal shear under plane strain condition. A stronger reduction in cross-sectional area due to necking is observed. Surface examination revealed conjugated shear bands at the same angles as the fracture surface with a width of 200–300 µm. A 1 mm TEM disc was punched out in one of the shear bands of specimen 3 as shown in Fig. 5c. A series of bright field and diffraction patterns showing the microstructural changes are given in Fig. 14. The bright field image in the centre shows that there is an elongated region where the microstructure is clearly different from the surrounding 20 nm scaled nanostructure. A diffraction pattern taken from an area that includes both the types of
Fig. 13. SEM micrograph of the surface of a fractured 3 mm nc Ni specimen tested at a low strain rate. Only one type of morphology can be observed (inset).
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Fig. 14. TEM images and SAD patterns of a locally deformed area taken from the shear band of a 20 mm nc sample deformed at 1.7 × 103s⫺1. (a) The central region in the shear band, (b) region from which the SAD pattern (c) is taken, (d) magnified region from which the SAD pattern (e) is taken.
microstructures (Fig. 14b,c) shows the typical diffraction rings occurring for a large number of grains of random orientations. Conversely, a diffraction pattern taken from a magnified view of the central area shows well-separated diffraction spots, which indicate the presence of limited number of grains. Tilting the sample revealed the presence of a 200-nm-sized grain in the region indicated by the smallest square (Fig. 14d). The black double-lined diffraction contrast relates to bending contours and indicates that the region covered by the double line is a single grain region, while the interruptions of it indicate that the grain is locally distorted. No dislocations have been detected within the larger grains, and often these larger grains seem to be free of internal stresses. The presence of these large grains have not been observed in undeformed specimens and can, therefore, be related to a localised deformation mode in high-speed deformed nc Ni.
5. Discussion and conclusions The microstructure and mechanical behaviour of the two batches of commercially available ED Ni were compared. It appears that although both the batches have the same mean grain size, as determ-
ined by XRD and TEM, they exhibit different microstructures and impurity contents. There is a clear difference in the concentration of the chemical impurities, especially in elements such as Cu, Al, W, Co and S. For most of the impurities, it is very difficult to quote their importance in the mechanical behaviour, except for S, which is known to be a critical impurity regarding embrittlement [28]. Batch B, which also contains a larger total amount of impurities, has a more pronounced {200} texture. The {200} structures can be related to fibre-like structures observed in TEM. The difference in microstructure is reflected in differences in hardness, yield stress and UTS. The hardness values are, however, contained within the prediction of the Hall–Petch slope, when experimental errors are taken into account. Differences in strength and ductility measured on 3 and 20 mm tensile specimens, with gauge lengths of 1.7 and 6 mm, respectively, are observed for nc Ni. A similar observation was made by Weertman et al. [8] for nc Cu. Experiments performed with single crystalline Ni revealed no difference in the mechanical properties between the two geometries, demonstrating that the explanation has to be found in microstructural effects. It was observed that the fractured surfaces of the small samples show a rough and irregular structure,
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which might originate from intergranular fracture as observed in the TEM, whereas the fractured surfaces of the bigger samples show a mixture of the rough structure together with a columnar one. This indicates that two types of microstructures at these length scales play a role in the deformation. The results suggest that a small gauge section could allow the participation of only one microstructure in the deformation process, whereas in bigger gauge sections, both type of microstructure would be involved. In addition, it could be invoked that the larger the surface to volume ratio in the gauge length, the more important the role of microvoids and microcracks in the overall deformation behaviour. This implies that sample size effects do contribute to the large scatter in the data of mechanical properties reported in literature. The amount of plastic strain that can be obtained in ED Ni is generally low and usually does not exceed 3%. A lack in plasticity at room temperature is observed in many, if not most, of the nc metals synthesised with the now existing methods [7]. Often grain boundary contamination and/or voids are mentioned as possible explanation. ED Ni contains a certain amount of S coming from the bath additives. Sulphur is known to have a strong tendency to segregate to more open regions such as cavities, grain boundaries and free surfaces, where the local concentration of sulphur can be orders of magnitude higher than the bulk value [28]. In coarse-grained Ni, sulphur contents of several hundred ppm are enough to cause grain boundary embrittlement [28]. Due to the considerable higher amount of interface surface, it is, however, not clear that a similar content of S in nc Ni causes embrittlement. Ebrahimi et al. [14] claim that using an additive-free sulphamate bath in the electrodeposition can reduce the S content. They measure, however, a maximal plastic strain of 1.3% for a mean grain size of 44 nm, which is not an improvement compared to the commercially available ED Ni. The present study, however, also demonstrates that there might be another reason for the lack in room temperature plasticity, i.e. the presence of cavities or pores along the grain boundaries and probably also inside some grains. These features are believed to generate during hydrogen codeposition [29] and are probably as important as
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the S content in provoking a limited plastic deformation [14]. ED nc Ni and polycrystalline Ni samples with the same geometry (20 mm total length) have been deformed at strain rates ranging from 10⫺5 to 103 s⫺1. Below 10 s⫺1, the UTS is about 4.5 times higher and the strain-rate sensitivity is one order of magnitude higher in the nc Ni. In both the cases, the work hardening rate increases with increasing strain rate. But in the nc samples, the work hardening rate decreases very rapidly during ongoing deformation, reaching the instability value of 1 below 2% plastic deformation. Lower work hardening values were also reported for nc copper compared to their coarse-grained counterparts tested over a large range of strain rates [30], however, the increase in strain-rate sensitivity was less pronounced. At higher strain rates, however, the increase in UTS is much more pronounced in the nc metal when strain rate increases by eight orders of magnitude. Together with the sudden increase in UTS, a change in fracture characteristics is observed. At the higher strain rates, the fracture surface is at 55– 65°, which corresponds to the angle of maximal shear under plain strain condition. There is also a stronger evidence of necking. Shear bands are observed on the surface in the deformed area, providing evidence for local deformation, as is reported in Refs. [31,32]. Bright field images from these areas reveal the presence of bigger grains up to 200 nm in the shear bands, while no evidence of dislocation activity is found. Kinetic energy originating from the high-speed deformation is probably responsible for local heat production, causing grain growth. It is known that ED Ni exhibits instantaneous grain growth already at 300 °C [16,20,33], temperatures, which can be expected inside a shear band [34]. In summary, the present study shows that commercially available ED nc Ni is not provided in a unique microstructure. Although grain sizes are similar within measurement errors, there are nonnegligible differences in texture and chemical impurities that result in non-negligible differences in yield strength, UTS and plastic strain. Compared to the coarse-grained Ni, nc Ni tested in tensile conditions shows a more pronounced increase in
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UTS with increasing strain rate, a higher strain-rate sensitivity for all applied strain rates and a lower work hardening. Fractured surfaces reveal two morphologies, only one of them leading to plasticity. At the highest strain rates, shear banding with local grain growth is observed. A size-dependent mechanical response is observed, which probably has to be related to the microstructural features and volume to surface ratio of the specimens. Acknowledgements The authors thank R. Schaeublin for many fruitful discussions, S. Van Petegem and co-workers for the Positron Annihilation measurements and EMPA for the assistance on the high-speed tensile machine. The NSF-project 21-52451.97 supported this work. References [1] Van Swygenhoven H, Derlet PM. Phys Rev B 2001;64:224. [2] Van Swygenhoven H, Spaczer M, Caro A. Acta mater 1999;47:3117. [3] Yamakov V, Wolf D, Salazar M, Phillpot SR, Gleiter H. Acta mater 2001;49:2713. [4] McFadden SX, Sergueeva AV, Krumel T, Martin JL, Mukherjee AK. Mater Res Soc Symp Proc 2001;634:B1. [5] Youngdahl CJ, Hugo RC, Kung H, Weertman JR. Mater Res Soc Symp Proc 2001;634:B1. [6] Derlet PM, Van Swygenhoven H. Phil Mag A 2002;82:1. [7] Koch CC, Morris DG, Lu K, Innoue A. MRS Bull 1999;24:54. [8] Weertman JR, Farkas D, Hemker KJ, Kung H, Mayo M, Mitra R, Van Swygenhoven H. MRS Bull 1999;24:44. [9] McFadden SX, Mishra RS, Valiev RZ, Zhilyaev AP, Mukherjee AK. Nature 1999;398:684. [10] Valiev RZ, Islamgaliev RK, Alexandrov IV. Prog Mater Sci 2000;45:103.
[11] Xiao C, Mirshams RA, Whang SH, Yin WM. Mater Sci Eng 2001;A301:35. [12] El-Sherik AM, Erb U, Palumbo G, Aust KT. Scr Metall Mater 1992;27:1185. [13] Natter H, Schmelzer M, Hempelmann R. J Mater Res 1998;13:1186. [14] Ebrahimi F, Bourne GR, Kelly MS, Matthews TE. Nanostruct Mater 1999;11:343. [15] Legros M, Elliott BR, Rittner MN, Weertman JR, Hemker KJ. Phil Mag A 2000;80:1017. [16] Dalla Torre F, Van Swygenhoven H, Victoria M, Scha¨ ublin R, Wagner W. Mater Res Soc Symp Proc 2001;634:B2. [17] Schiøtz J, Vegge T, Di-Tolla FD, Jacobsen KW. Phys Rev B 1999;60:11. [18] Krill CE, Birringer R. Phil Mag A 1998;77:621. [19] Bailat C. PhD dissertation. The`se No. 2011, Lausanne: EPFL; 1999. [20] Wang N, Wang Z, Aust KT, Erb U. Acta Metall Mater 1995;45:1655. [21] Ru¨ hle M, Wilkens M. Cryst Latt Def 1975;6:129. [22] Van Petegem S. Private communication. [23] Hughes GD, Smith SD, Pande CS, Johnson HR, Armstrong RW. Scr Metall Mater 1986;20:93. [24] El-Sherik AM, Erb U, Palumbo G, Aust KT. Scr Metall Mater 1992;27:1185. [25] Sevillano JG. In: Mugrabi H, editor. Plastic deformation and fracture of materials, 6. Weinheim: Verlag Chemie; 1993, p. 19. [26] Kolsky H. Proc Res Soc Lond 1949;B62:676. [27] Follansbee PS, Huang JC, Gray GT. Acta Metall Mater 1990;38:1241. [28] Seah MP, Leah C. Phil Mag 1975;31:627. [29] Merchant HD. Proceedings of TMS Meeting of Defect Structure, Morphology and Properties of Deposits. The Minerals, Metals & Materials Society. 1995, p. 1. [30] Jia D, Ramesh KT, Ma E, Lu L, Lu K. Scr Mater 2001;45:613. [31] Jia D, Wang M, Ramesh KT, Ma E, Zhu YT, Valiev RZ. Appl Phys Lett 2001;79:611. [32] Jia D, Ramesh KT, Ma E. Scr Mater 2000;42:73. [33] McFadden SX, Zhilyaev AP, Mishra RS, Mukherjee AK. Mater Lett 2000;45:345. [34] Andrade U, Meyers MA, Vecchio KS, Chokshi AH. Acta Metall Mater 1994;42:3183. [35] Van Swygenhoven, H., Derlet, PM, Hasnaoui, A Phys Rev B 2002, in press.
Nanocrystalline electrodeposited Ni: microstructure and tensile properties F. Dalla Torre a, H. Van Swygenhoven a,∗, M. Victoria b b
a Paul Scherrer Institute, CH-5232 Villigen, Switzerland CRPP-Fusion Technology Materials, EPFL, CH-5232 Villigen, Switzerland
Received 2 April 2002; accepted 15 May 2002
Abstract The microstructure of commercially available nanocrystalline (nc) electroplated Ni foils is studied by means of Xray diffraction and transmission electron microscopy. It is shown that the microstructure is inhomogeneous and batchdependent. Tensile properties at strain rates between 10⫺5 and 103 s⫺1 are studied and compared with the results of coarse-grained Ni. Data on strength, strain-rate sensitivity and work hardening are presented. At the highest strain rates, shear banding with local grain growth is observed in the nc structure. It is also suggested that the differences found in nc Ni for 3 and 20 mm tensile specimens are the size effects related to the inhomogeneous microstructure. 2002 Published by Elsevier Science Ltd on behalf of Acta Materialia Inc. Keywords: Nanocrystalline; Mechanical properties; Microstructure; Nickel; Shear bands
1. Introduction It is well known that in polycrystalline metals, an increase in hardness and strength can be obtained by reducing the grain size to the nanometer scale. There is, however, no clear understanding of the deformation mechanisms that operate at these grain sizes. It is expected that plasticity is, to a lesser extent, driven by dislocation mechanisms and that deformation will be accommodated in the grain boundaries. Computer simulations in full 3D nanostructures [1,2] as well as in columnar ∗ Corresponding author. Paul Scherrer Institute, Building WHGA/343, CH-5232 Villigen PSI, Switzerland. Tel.: +41-56310-21-11; fax: +41-56-310-21-99. E-mail address: [email protected] (H. Van Swygenhoven).
nanostructures [3] suggested grain boundary sliding as the main deformation process [1]. However, direct experimental observation of deformation mechanism at atomic scale is still missing. In situ deformation studies using transmission electron microscopes (TEM) [4,5] do not provide convincing results since possible relaxations in the nanosized grain boundaries due to the thinning procedure tend to obscure interpretations [6]. Another concern, and major obstacle, in the investigation of the deformation mechanism is the limited ductility reported in these nanostructures [7], which are believed to be mainly provoked by pre-existing flaws, such as impurities and porosity [8], introduced during the fabrication of the material. Therefore, a lot of attention is paid to electrodeposited (ED) and severe plastic deformed
1359-6454/02/$22.00 2002 Published by Elsevier Science Ltd on behalf of Acta Materialia Inc. PII: S 1 3 5 9 - 6 4 5 4 ( 0 2 ) 0 0 1 9 8 - 2
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nanocrystalline (nc) samples [9–12] as it is believed that these techniques can produce porosity-free samples. In recent years, efforts have been made to relate the nanostructure to the deposition parameters in order to obtain texture-free nanometer-sized grain structures with high chemical purity [13,14]. The preferential growth of crystals and the resulting texture are indeed common problems in the electrodeposition of metals and are difficult to control. Therefore, reports on mechanical properties have to be accompanied by an accurate microstructural investigation. Usually grain size and microstrain are provided by means of X-ray diffraction (XRD), but often little information is available on TEM observations [11–13]. Another major difficulty involved in performing mechanical tests on nc metals is often the limited amount of material provided by the above-mentioned techniques. Electrodeposition does not suffer from this restriction, although the sample thickness usually does not exceed 0.3 mm. As other methods, such as inert gas condensation and ball milling, do provide only a limited amount of material, miniaturised tensile machines have been developed allowing the comparison of the mechanical properties of samples synthesised by different methods [15,16]. Recent computer simulation results [1,17,35] have provided an analysis of the deformation mechanism. These simulations, though, are performed at very high strain rates. Part of the present experimental effort includes testing at high strain rates, up to 103 s⫺1, in order to extrapolate to the ultra-high strain rates reached in the simulations. This article provides a detailed description of the microstructure of a commercially available ED Ni and its room temperature mechanical behaviour under uniaxial tension. The microstructure of the ED Ni samples is characterised by means of their mean grain size, internal strains and texture deduced from XRD and TEM. Different batches are compared in order to investigate the reproducibility of the electrodeposition technique. Tensile tests are performed at strain rates ranging from 10⫺5 to 103 s⫺1 and are compared with the behaviour of coarse-grained Ni.
2. Experimental procedures Two batches (A and B) of electroplated nc Ni foils with a purity of 99.95% and a cold rolled coarse-grained Ni foil (purity 99.98%) were purchased from Goodfellow Metals. The thickness of the foils is 200–250 µm for the nc Ni and 1 mm for the coarse-grained Ni. Tensile specimen of coarsegrained Ni were cut along the rolling direction, mechanically polished to a thickness of approximately 0.5 mm and afterwards annealed at 700 °C. Chemical analysis on both the nc Ni batches was made via inductively coupled plasma optical emission spectrometry (ICP-OES). The density of the batches A and B was determined by the Archimedes immersion method using a Mettler microgram balance and diethyl-phthalate as a reference liquid. For both the batches, a density of 99.5 ± 0.5% was obtained. XRD measurements were performed with a Siemens D500 X-ray detector. Average grain size and lattice strain were calculated using the Warren– Averbach approach after Krill and Birringer [18] for the {111} / {222} and {200} / {400} family planes. Texture was measured using the relative intensities of the {111} and {200} peaks normalised over a standard powder diffraction pattern of (100% random) Ni powder. Conventional TEM investigations were made with a Jeol 2010 operated at 200 kV. TEM specimens were electropolished with a TENUPOL 5 of Struers. A solution of 3% perchloric acid, 30% ethylenglycol and 67% methanol at 15 V and ⫺10 °C was used. Grain boundaries of approximately 500 grains were delineated manually with adobe illustrator software on digitised TEM bright field images to determine the grain size distribution, which is fitted with a lognormal function. Because of the high lattice distortion of the microstructure, bright field images were preferred over the dark field images for the measurement of grain size. The high lattice distortion in the grains leads to a blurred diffraction contrast resulting in smaller grain sizes when measured by dark field images. Tensile tests were performed at room temperature with different tensile machines: (1) a strain rate controlled miniaturised tensile machine using 3 mm long dog-bone samples [16] at a strain rate
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of 2 × 10⫺4s⫺1; (2) two conventional tensile testing machines (Zwick and Schenk) using 20 mm dogbone shaped specimens at strain rates between 10⫺5 and 10⫺1 s⫺1 [19]; and (3) a Roell Amsler, type 1852, servo-hydraulic high-speed tensile machine designed for dynamic tension and compression tests with a maximal cross-head speed of 12 m s⫺1. When the elastic tensile slope is compared with the one measured in the Zwick or Schenk machine, it appears that the stiffness of this machine is 2.3 times lower. Therefore, strain values of the Roell Amsler machine were divided by this factor, so that strains can be compared. The yield strength and the plastic strain are calculated using the intercept of the stress–strain curve with the 0.2% total strain offset parallel to the elastic slope. The tensile specimens, shown in Fig. 5, have been machined by a wire Electro-DischargingMachine (EDM). The 3 mm long tensile specimens have a gauge length of 1.72 mm, thickness, approximately 0.20 mm and width, approximately 0.25 mm. The 20 mm tensile samples have a gauge cross-section of 2.5 × 0.2 mm and a gauge length of 6 mm. Local heat generated by the 0.1 mm thick wire during the spark erosion caused local grain growth to a 10 µm thick layer at the surfaces. Therefore, the specimens were additionally electropolished during 30 s using the same conditions as used for producing the TEM discs to remove the altered surface. Microhardness measurements were performed with a Vickers pyramid indenter, installed on a Zeiss Jenaphot 2000 optical microscope and were monitored by a computer. Mean hardness values are calculated from 4 to 9 µm indentation depths. Fractured surfaces of tensile specimens and indentations were examined by scanning electron microscopy using a Topcon instrument.
3. Microstructure 3.1. Chemical composition Table 1 shows the chemical analysis of the main impurities (ppm) for batches A and B. Metallic impurities in ED samples usually originate from
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the impurities in the anode material and from the chemicals used for the electrolytic bath. Cobalt and iron, for instance, are common impurities in nickel sulphate, sulphur originates from bath additives, such as natrium-saccharin (Na-C7H4NO3S) [13]. Light element impurities such as H, N and O could not be measured, but are expected to be present in the material as well. Co-deposition of hydrogen generated by the dissociation of water as well as oxidation processes in the solution is a common chemical by-product during electrodeposition. Molecules of additives, such as Na-Saccharin often used as grain-growth inhibitors may be adsorbed by the surface. Natter et al. [13] detected 110, 620 and 300 ppm hydrogen, nitrogen and oxygen, respectively, in a 17 nm grain size sample. The comparison of the amount of impurities between the batches A and B shows significant difference for some elements. The Cu content in batch B is almost three times higher. Also Al, Co and W are present in higher amounts in batch B compared to batch A. The sulphur content of batch A is, however, 15% higher. Also higher concentration of Ca and Pb are found in batch A. The differences in impurity content, indicating the difficulty in reproducing identical compositions, could play a non-negligible role in the material properties. 3.2. Texture, grain size and grain size distribution Table 2 summarises the microstructural observations of texture, grain size and strain. Compared to a standard Ni sample with a fully random crystal orientation, both the batches show preferential orientation along the {200} family planes but to different degrees. Batch A shows a 50% random orientation whereas batch B shows 30% random orientation. Data obtained from the solution side and the substrate side of both the batches give similar results. Polishing the samples results in a slightly higher observed {200} texture. TEM observations also reveal a textured microstructure. Fig. 1 shows TEM bright field images and the corresponding selected area diffraction (SAD) patterns for batch A (left) and batch B (right). The closest first ring of the diffraction pattern corresponds to the {111} peak and the second
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Table 1 Chemical impurity content of the batches A and B Element
Batch A (ppm)
Batch B (ppm)
6.5 8.75 881 155 127 19 883 ⬍
45.9 2.56 104 473 135 ⬍ 746 263
Al Ca Co Cu Fe Pb S W
Detection limit (ppm) 1 0.02 0.5 0.1 0.1 10 5 10
Table 2 Summary of the obtained microstructural measurements on grain size, strain and texture using XRD and TEM XRD {111}/ {222}
Batch A Batch B
TEM
Grain size (nm)
Microstrain (%)
Texture
Mean grain size (nm)
Peak value (nm)
Peak width
20.8 ± 3.0 21.4 ± 4.2
0.403 ± 0.015 0.393 ± 0.070
50% ± 6 30% ± 4
21.3 ± 7.8 20.6 ± 8.6
17.9 16.1
0.36 0.38
Fig. 1. TEM bright field image of batches A (left) and B (right) together with the corresponding SAD patterns (750 nm ring diameter).
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to the {200} peak of the Ni fcc structure. The difference in the relative intensities of the two rings in batches A and B demonstrates the stronger {200} texture for batch B. No additional peaks could be measured indicating that no second phase is present in the ED Ni foils. Mean grain sizes and microstrain measured with XRD reveal similar values for batches A and B. When using the {111} / {200} reflections, the corresponding values for batch A are a mean size of 20.8 ± 3.0nm and a microstrain of 0.40% whereas for batch B they are 21.4 ± 3.0nm with 0.39% strain. When the {200} / {400} reflections are used, again no notable difference between A and B is found, but here, due to the low {400} peak intensity, a large scattering between different measurements is observed (12–60 nm for the mean grain size and 0.54–0.98% for the microstrain). Grain size calculated from bright field TEM images resulted in similar mean grain sizes: 21.3 ± 7.8nm for batch A and 20.6 ± 7.8nm for batch B. Fig. 2 shows the grain size distributions for both the batches fitted with a lognormal distribution function. For batch A (B), a peak value of 17.9 nm (16.1 nm) and a peak width of 0.36 (0.38) are obtained. The ratios between the maximal and the minimal grain diameters are 1.60 for batch A and 1.63 for batch B. 3.3. Microstructural investigation by TEM TEM bright field images of both the batches (Fig. 1) show regions where grain shapes and grain
Fig. 2.
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boundaries can be clearly observed along with other regions of high density of diffraction contrast spatial variation. It is assumed that these contrasts originate mainly from microstrains, as suggested by XRD. This makes grain boundary features difficult to distinguish from other contrasts, such as bending contours and thickness fringes. Moreover, groups of grains possessing only small misorientation angles can be observed, as is shown in the higher magnification bright field image in Fig. 3a. Grains can be distinguished by their change in contrast at the grain boundaries. However, some grain boundaries show weak contrast. Tilting experiments demonstrate the presence of groups of grains with small misorientation angles. The Moire´ fringes observed in Fig. 3a can probably be ascribed to the superposition of such grains. A similar observation was made by Wang et al. [20], where the authors further suggest that these regions of small misfit are the nucleation sites of abnormal grain growth. The observation of the microstructure at high magnification revealed no clear evidence for the presence of dislocations. It should be noted that their contrasts could be convoluted with the contrast originated from microstrain in such a way that their identification is impeded. The stronger {200} texture, observed in batch B, can be associated with fibre-like structures as shown in the bright field image and the corresponding SAD pattern in Fig. 3b. The smeared-out {200} diffraction spots correspond to the dark
Grain size distributions for batches A (a) and B (b) using TEM bright field pictures fitted to a lognormal distribution.
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Fig. 3. TEM micrographs of batch B at higher magnifications: (a) region where grains can be distinguished by the change of contrast at grain boundaries. Moire´ fringes are due to the superposition of grains lying on the top of each other; (b) region with fibre-like texture. The inset shows the SAD pattern (300 nm) indicating the {200} texture.
wavy fibre-like structures in the bright field image. The fibre-like structures are inhomogeneously distributed with distances typically from 100 to 1000 nm between the groups. Fibre-type structures have also been observed in Ref. [12]. The TEM defocusing technique [21] revealed the presence of small voids with sizes between 1 and 2 nm, often situated close to or at the grain boundaries, as shown in Fig. 4 (taken from batch B). Their origin lies probably in hydrogen being incorporated during electrodeposition. The presence of voids or cavities is also confirmed by positron lifetime measurements, where the lifetimes of 160 and 350 ps are found, which relate to single vacancies and voids that are approximately 1 nm in size, respectively [22].
4. Mechanical behaviour Fig. 5a–c shows photographs of typical tensile specimens used in this study. Fig. 5b shows a 3 mm sample for batches A and B before and after the deformation at 10⫺4 s⫺1 strain rate. Examples of the 20 mm samples are shown in Fig. 5c: before testing (1), deformed at 10⫺4 s⫺1 (2), and deformed
Fig. 4. TEM image of batch B showing pores of 1–2 nm in size. The inset shows pore concentrations in the grain interior.
at 103 s⫺1 (3). In Section 4.1, the differences in mechanical behaviour between the batches are discussed. The influence of samples size of batch B is discussed in Section 4.2, and the effect of strain rate is discussed in Section 4.3.
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for the yield strength and 2.0 GPa (1.8 GPa) for the ultimate tensile strength (UTS). Both the batches showed post-necking elongation. Microhardness tests showed a similar trend as observed for the yield strength in tension—a lower hardness for batch B. The mean hardness values are 6.0 ± 0.2 and 5.5 ± 0.5GPa for batches A and B, respectively. Fig. 6 shows a Hall–Petch plot, representing our results for batches A (circles) and B (diamond) as well as the values taken from the literature [23,24]. The grain sizes are determined by XRD (full datapoints) and by the peak value of the lognormal distribution function obtained by TEM (open datapoints). 4.2. Influence of specimen size Fig. 5. Tensile specimens. (a) The size difference between big and miniaturised specimens. (b) The 3 mm specimens before testing and deformation at 10⫺4 s⫺1 for batches A and B. (c) The 20 mm specimens before testing (1) and after deformation at 10⫺4 s⫺1 (2), and at 103 s⫺1 (3). The arrow indicates the shear band, which has an orientation of about 60° with respect to the tensile axis.
Comparative tests were performed to investigate the influence of sample size in tensile testing, as
4.1. Differences in tensile properties between batches A and B Table 3 summarises the results obtained from the tensile tests, nanoindentation and microhardness measurements for both the batches. The tensile tests were performed on the 3 mm samples at a strain rate of 2 × 10⫺4s⫺1. Higher ductility and yield strength were measured for batch A. The following data are obtained for batch A (batch B): 4% (2.6%) for the plastic strain, 1.6 GPa (1.4 GPa)
Fig. 6. Hall–Petch plot representing our results among data found in literature.
Table 3 Summary of the measured mechanical properties (results of 3 mm tensile specimens) Tensile tests UTS (GPa) Batch 2.04 ± 0.10 A Batch 1.82 ± 0.09 B
Microhardness sy (GPa)
Plastic strain (%)
Hv (GPa)
1.60 ± 0.10
4.0 ± 0.2
6.0 ± 0.3
1.43 ± 0.09
2.6 ± 0.6
5.5 ± 0.1
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samples revealed the presence of necking over the whole cross-section and in most cases a homogenous rough morphology at the fractured surface indicating ductile fracture. The 20 mm samples, however, show the presence of two types of morphologies, the same rough structure and a columnar structure (see Section 3.3). It suggests that these different morphologies also represent a difference in the crystallographic orientation and/or in the shape of grains. Whether the columnar structure can be related to the fibre-type texture observed by TEM is a matter of further research. Fig. 7. Tensile curves for 3 and 20 mm specimens of batch B. Yield and UTS are lower for bigger specimens.
shown in the stress–strain plot in Fig. 7. Table 4 shows the UTS obtained for different strain rates of six tensile tests carrried out for the 3 and the 20 mm specimens. Mean values of 1849 ± 51 and 1388 ± 61 MPa were obtained for the two specimen sizes. For single crystal Ni, however, no difference was measured between the two types of samples. This means that the reason for this apparent size effect has to be found in microstructural effects, such as the presence of surface imperfections and/or the inhomogeneity of the microstructure in the gauge length. Careful examination of the EDM-cut surfaces after polishing revealed occasionally the presence of very small microcracks. Accumulation of the effect of these features is sample size dependent and could contribute to an overall lower deformation stress. On the other hand, examination of deformed 3 mm
4.3. Influence of strain rate Fig. 8 shows UTS as a function of strain rate for both nc and coarse-grained Ni. Instead of plotting the stress at a constant strain (as it is usually done for similar plots), the UTS was favoured since only few data points of stress at a certain amount of strain show overlap for both types of materials. The strain-rate sensitivity m, defined as m ⫽ (∂lns / ∂lne˙ )e, where s is the true stress, e˙ , the strain rate and e, the true strain, which is, however, calculated between 2 and 3% strain for nc and between 2 and 20% strain for coarse-grained Ni. The condition for stable deformation can be formulated [25] in terms of the strain-rate sensitivity and the work hardening rate Θ defined as Θ ⫽
Table 4 Measurements of the UTS for 3 and 20 mm specimens 3 mm tensile specimens 20 mm tensile specimens UTS Strain rate (s⫺1) (MPa)
UTS (MPa)
Strain rate (s⫺1)
7 × 10⫺5 2 × 10⫺4 2 × 10⫺4 2 × 10⫺3 1 × 10⫺2 1 × 10⫺2
1448 1397 1278 1396 1437 1374
5.5 5.5 5.5 5.5 5.5 5.5
1823 1828 1818 1921 1905 1800
× 10⫺5 × 10⫺4 × 10⫺4 × 10⫺3 × 10⫺3 × 10⫺2
Fig. 8. UTS versus the logarithm of strain rate for nc and coarse-grained Ni. The grey bar around the highest strain rate data points of coarse-grained Ni marks the uncertain data.
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1 (∂s / ∂e)e˙ . For materials with low strain-rate sensiσ tivity, plastic instability is reached when Θⱕ1. 4.3.1. Low strain rates Fig. 9a,b shows true stress–true strain curves and the work hardening rate Θ measured on 20 mm specimens of batch B using four different strain rates in the lower strain rate regime (10⫺5–10⫺2 s⫺1). Symbols are used in this figure to distinguish between curves and do not represent specific data points. As can be seen in Fig. 9, the maximal strains that can be obtained, decrease with increasing strain rate, while the UTS remains approximately constant. The UTS is, however, up to 4.5
Fig. 9. (a) True stress–true strain curve for nc Ni at strain rates in the lower strain rate regime. (b) Work hardening rate ⍜ versus true stress for the different strain rates.
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times higher (Fig. 8) in nc Ni. Plastic strain ranges from 2.2% for the lowest strain rate to 0.4% for the highest. The strain-rate sensitivity m is one order of magnitude higher for the nc (0.01–0.03) as compared to the coarse-grained Ni (0.001–0.004) for strain values between 2 and 20% total strain. The hardening rate (Fig. 9b) increases when strain rate increases from 5.5 ± 10⫺2 to 5.5 ± 10⫺5s⫺1. When compared with the coarse-grained material, a value of ⌰ ⫽ 10 at 1360 MPa is obtained from the ED Ni whereas ⌰ ⫽ 14 at 130 MPa is obtained for the coarse-grained Ni. Both represent values at 2% plastic strain. 4.3.2. High strain rates At higher strain rates the UTS increases for both the materials, the increase in the nc material is, however, more pronounced, up to 70% (2500 MPa) when strain rate is increased to 103 s⫺1. The strainrate sensitivity at these high stress levels is 0.47. No value for m could be calculated for the coarsegrained Ni, since the data at high strain rates (above 4 × 102s⫺1) have to be taken with caution, as discussed subsequently. Similar to the Kolsky bar technique [26], where a bullet is shot to a rod that propagates the compressive stress wave onto the specimen, the highspeed tensile machine used in this study transfers the kinetic energy of the accelerated plunger onto the specimen by an instantaneous applied force that propagates as a tensile stress wave in the specimen load train. With increasing loading velocity, the appearance of one or more unloading regions is observed during the tests (see tensile curve at 8.1 × 102s⫺1 in Fig. 10). This is interpreted as the arrival of the first reflection of the tensile wave on the specimen, which is now a compressive (unloading) stress wave. This wave superimposes on the initial tensile wave, resulting in a non-trivial wave shape. Fourier transformation analysis of the tensile data showed two main oscillations with frequencies of approximately 5.4 and 6.4 Hz and waves with these frequencies are seen to continue to propagate after the failure of the specimen. Using low pass filtering to skip the undesired higher frequencies, a continuous straight loading curve can be extracted, but the stress amplitude value is also affected. The total strain on the other
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Fig. 10. Engineering stress–strain curves at high strain rates for coarse-grained Ni.
Fig. 11. Engineering stress–strain curves at high strain rates for nc Ni.
hand, could be measured on the fractured specimens and this corresponds well with the data given by the strain gauges in the machine. In other words, the UTS at strain rates above 4 × 102s⫺1 are to be taken with caution and represent only trends. The data for nc Ni are, however, correct, since all the samples broke before the tensile wave was reflected. Fig. 10 shows a plot of the engineering stress– strain curves of coarse-grained Ni for high strain and rates ranging between 5.4 × 10⫺1 2 ⫺1 8.1 × 10 s . The maximal observed true stress value is 0.47 GPa and the maximal measured plastic strain is 17% of which 2–3% contributes to post-necking elongation. Higher strain rates result in an increase in strength. The effect on the work hardening is, however, less pronounced for strain rates between 5 × 101 and 1.3 × 102s⫺1. A more pronounced increase in ⍜ for higher strain rates can be expected in coarse-grained Ni [27]. All tested coarse-grained Ni samples show uniform deformation over the whole gauge length before the necking sets in and leads to failure. At a strain rate of 8 × 102s⫺1, oscillations are observed in the stress–strain curve, representing the reflections of the tensile wave. Fig. 11 shows the engineering stress–strain curves at high strain rates for nc Ni, measured on the 20 mm samples. Plastic strains range between 1.2 and 1.8%, but there is, however, no clear relation with the applied strain rate. At the highest strain rates, the yield point is less pronounced,
however, stress values of 2.5 GPa are supported by the nanostructure. The work hardening rate seems to increase with increasing strain rate, the values are, however, difficult to compare since plastic instability sets in very early. Values of ⌰ ⫽ 1 are reached already before 1% plastic strain. 4.4. Microstructural observation of fracture surfaces and deformed areas 4.4.1. Low strain rates The 20 mm specimens have a fractured surface perpendicular to the tensile direction and this is true for all the strain rates up to 102 s⫺1 (Fig. 5c). Minor sign of plasticity could be detected on the flat specimen surface along with the little evidence of necking. An SEM picture of some details of the fractured surface is given in Fig. 12. Within the same fractured surface, two different structures could be observed, marked by the letters R and C. The zone R represents a very rough and irregular surface structure, whereas zone C shows columnar fractured lines perpendicular to the foil plane. In this picture, the two regions are relatively small sized, but the dimensions of zones R and C can vary quite strongly, both reaching till 50% of the fractured surface. SEM examinations showed clearly that only the structure responsible for the rough fractured surface contributed to the observed plasticity. In samples, where a large percentage of the fractured surface was of type R, local reduction in gauge cross-section was observed, indicating
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Fig. 12. SEM micrographs of the surface of a fractured 20 mm nc Ni specimen tested at a low strain rate. The fractured surface shows two types of morphologies indicated by the letters R (rough) and C (columnar).
‘local’ necking. Most of the fractured surfaces of the 3 mm samples, deformed at low strain rates, showed only the morphologies of type R, as can be seen in Fig. 13. As was mentioned before, those samples showed post-necking elongation. The difference in fracture mechanism is probably related to the inhomogeneous microstructure revealed in TEM. 4.4.2. High strain rates At higher strain rates, a clear change in fracture mode is observed in the 20 mm samples. The angle of the fracture surface with the tensile direction changed from ~90 to 65–55°, corresponding to the
angle of maximal shear under plane strain condition. A stronger reduction in cross-sectional area due to necking is observed. Surface examination revealed conjugated shear bands at the same angles as the fracture surface with a width of 200–300 µm. A 1 mm TEM disc was punched out in one of the shear bands of specimen 3 as shown in Fig. 5c. A series of bright field and diffraction patterns showing the microstructural changes are given in Fig. 14. The bright field image in the centre shows that there is an elongated region where the microstructure is clearly different from the surrounding 20 nm scaled nanostructure. A diffraction pattern taken from an area that includes both the types of
Fig. 13. SEM micrograph of the surface of a fractured 3 mm nc Ni specimen tested at a low strain rate. Only one type of morphology can be observed (inset).
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Fig. 14. TEM images and SAD patterns of a locally deformed area taken from the shear band of a 20 mm nc sample deformed at 1.7 × 103s⫺1. (a) The central region in the shear band, (b) region from which the SAD pattern (c) is taken, (d) magnified region from which the SAD pattern (e) is taken.
microstructures (Fig. 14b,c) shows the typical diffraction rings occurring for a large number of grains of random orientations. Conversely, a diffraction pattern taken from a magnified view of the central area shows well-separated diffraction spots, which indicate the presence of limited number of grains. Tilting the sample revealed the presence of a 200-nm-sized grain in the region indicated by the smallest square (Fig. 14d). The black double-lined diffraction contrast relates to bending contours and indicates that the region covered by the double line is a single grain region, while the interruptions of it indicate that the grain is locally distorted. No dislocations have been detected within the larger grains, and often these larger grains seem to be free of internal stresses. The presence of these large grains have not been observed in undeformed specimens and can, therefore, be related to a localised deformation mode in high-speed deformed nc Ni.
5. Discussion and conclusions The microstructure and mechanical behaviour of the two batches of commercially available ED Ni were compared. It appears that although both the batches have the same mean grain size, as determ-
ined by XRD and TEM, they exhibit different microstructures and impurity contents. There is a clear difference in the concentration of the chemical impurities, especially in elements such as Cu, Al, W, Co and S. For most of the impurities, it is very difficult to quote their importance in the mechanical behaviour, except for S, which is known to be a critical impurity regarding embrittlement [28]. Batch B, which also contains a larger total amount of impurities, has a more pronounced {200} texture. The {200} structures can be related to fibre-like structures observed in TEM. The difference in microstructure is reflected in differences in hardness, yield stress and UTS. The hardness values are, however, contained within the prediction of the Hall–Petch slope, when experimental errors are taken into account. Differences in strength and ductility measured on 3 and 20 mm tensile specimens, with gauge lengths of 1.7 and 6 mm, respectively, are observed for nc Ni. A similar observation was made by Weertman et al. [8] for nc Cu. Experiments performed with single crystalline Ni revealed no difference in the mechanical properties between the two geometries, demonstrating that the explanation has to be found in microstructural effects. It was observed that the fractured surfaces of the small samples show a rough and irregular structure,
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which might originate from intergranular fracture as observed in the TEM, whereas the fractured surfaces of the bigger samples show a mixture of the rough structure together with a columnar one. This indicates that two types of microstructures at these length scales play a role in the deformation. The results suggest that a small gauge section could allow the participation of only one microstructure in the deformation process, whereas in bigger gauge sections, both type of microstructure would be involved. In addition, it could be invoked that the larger the surface to volume ratio in the gauge length, the more important the role of microvoids and microcracks in the overall deformation behaviour. This implies that sample size effects do contribute to the large scatter in the data of mechanical properties reported in literature. The amount of plastic strain that can be obtained in ED Ni is generally low and usually does not exceed 3%. A lack in plasticity at room temperature is observed in many, if not most, of the nc metals synthesised with the now existing methods [7]. Often grain boundary contamination and/or voids are mentioned as possible explanation. ED Ni contains a certain amount of S coming from the bath additives. Sulphur is known to have a strong tendency to segregate to more open regions such as cavities, grain boundaries and free surfaces, where the local concentration of sulphur can be orders of magnitude higher than the bulk value [28]. In coarse-grained Ni, sulphur contents of several hundred ppm are enough to cause grain boundary embrittlement [28]. Due to the considerable higher amount of interface surface, it is, however, not clear that a similar content of S in nc Ni causes embrittlement. Ebrahimi et al. [14] claim that using an additive-free sulphamate bath in the electrodeposition can reduce the S content. They measure, however, a maximal plastic strain of 1.3% for a mean grain size of 44 nm, which is not an improvement compared to the commercially available ED Ni. The present study, however, also demonstrates that there might be another reason for the lack in room temperature plasticity, i.e. the presence of cavities or pores along the grain boundaries and probably also inside some grains. These features are believed to generate during hydrogen codeposition [29] and are probably as important as
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the S content in provoking a limited plastic deformation [14]. ED nc Ni and polycrystalline Ni samples with the same geometry (20 mm total length) have been deformed at strain rates ranging from 10⫺5 to 103 s⫺1. Below 10 s⫺1, the UTS is about 4.5 times higher and the strain-rate sensitivity is one order of magnitude higher in the nc Ni. In both the cases, the work hardening rate increases with increasing strain rate. But in the nc samples, the work hardening rate decreases very rapidly during ongoing deformation, reaching the instability value of 1 below 2% plastic deformation. Lower work hardening values were also reported for nc copper compared to their coarse-grained counterparts tested over a large range of strain rates [30], however, the increase in strain-rate sensitivity was less pronounced. At higher strain rates, however, the increase in UTS is much more pronounced in the nc metal when strain rate increases by eight orders of magnitude. Together with the sudden increase in UTS, a change in fracture characteristics is observed. At the higher strain rates, the fracture surface is at 55– 65°, which corresponds to the angle of maximal shear under plain strain condition. There is also a stronger evidence of necking. Shear bands are observed on the surface in the deformed area, providing evidence for local deformation, as is reported in Refs. [31,32]. Bright field images from these areas reveal the presence of bigger grains up to 200 nm in the shear bands, while no evidence of dislocation activity is found. Kinetic energy originating from the high-speed deformation is probably responsible for local heat production, causing grain growth. It is known that ED Ni exhibits instantaneous grain growth already at 300 °C [16,20,33], temperatures, which can be expected inside a shear band [34]. In summary, the present study shows that commercially available ED nc Ni is not provided in a unique microstructure. Although grain sizes are similar within measurement errors, there are nonnegligible differences in texture and chemical impurities that result in non-negligible differences in yield strength, UTS and plastic strain. Compared to the coarse-grained Ni, nc Ni tested in tensile conditions shows a more pronounced increase in
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UTS with increasing strain rate, a higher strain-rate sensitivity for all applied strain rates and a lower work hardening. Fractured surfaces reveal two morphologies, only one of them leading to plasticity. At the highest strain rates, shear banding with local grain growth is observed. A size-dependent mechanical response is observed, which probably has to be related to the microstructural features and volume to surface ratio of the specimens. Acknowledgements The authors thank R. Schaeublin for many fruitful discussions, S. Van Petegem and co-workers for the Positron Annihilation measurements and EMPA for the assistance on the high-speed tensile machine. The NSF-project 21-52451.97 supported this work. References [1] Van Swygenhoven H, Derlet PM. Phys Rev B 2001;64:224. [2] Van Swygenhoven H, Spaczer M, Caro A. Acta mater 1999;47:3117. [3] Yamakov V, Wolf D, Salazar M, Phillpot SR, Gleiter H. Acta mater 2001;49:2713. [4] McFadden SX, Sergueeva AV, Krumel T, Martin JL, Mukherjee AK. Mater Res Soc Symp Proc 2001;634:B1. [5] Youngdahl CJ, Hugo RC, Kung H, Weertman JR. Mater Res Soc Symp Proc 2001;634:B1. [6] Derlet PM, Van Swygenhoven H. Phil Mag A 2002;82:1. [7] Koch CC, Morris DG, Lu K, Innoue A. MRS Bull 1999;24:54. [8] Weertman JR, Farkas D, Hemker KJ, Kung H, Mayo M, Mitra R, Van Swygenhoven H. MRS Bull 1999;24:44. [9] McFadden SX, Mishra RS, Valiev RZ, Zhilyaev AP, Mukherjee AK. Nature 1999;398:684. [10] Valiev RZ, Islamgaliev RK, Alexandrov IV. Prog Mater Sci 2000;45:103.
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