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«STRUCTURE AND PROPERTIES OF ELECTRODEPOSITED NANOCRYSTALLINE NI AND NI-FE ALLOY CONTINUOUS FOILS by Jason Derek Giallonardo A thesis submitted in ...»

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In order to assess the presence of edge dislocations, numerous grains having an [011] zone axis were examined from the series of HR-TEM images. Very few edge dislocations were observed which is consistent with earlier findings on electrodeposited nanocrystalline Ni, NiP, and Ni-Fe alloys [Mehta et al. (1995), Giallonardo et al. (2012)]. Wu et al. (2007) also examined HR-TEM images of as-received nanocrystalline Ni samples produced using the electrodeposition technique and very few dislocations were observed. Note that there is considerable strain contrast at some of the grain boundaries as will be discussed in later parts of this section.

Triple junctions, high and low angle grain boundaries, and twin boundaries were examined in detail. Fig. 4.16 shows a well defined triple junction from sample no. 5 (Ni, 23 nm). The zone axes for grains, A, B, and C, are [001], [112] and, [011], respectively. Since

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Figure 4.16.

HR-TEM image of a triple junction in sample no. 5 (Ni, 23 nm). To the right of the image are the FFT patterns for grains A, B, and C corresponding to zone axes [001], [112], and [011], respectively.

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Figure 4.17.

HR-TEM image of the boundary between grains A and B from Fig. 4.16. The respective zone axes for grains A and B are [001] and [112].

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the zone axes are quite different, the grain boundaries are considered to be high angle boundaries. Another section of the grain boundary between grains A and B away from the triple junction is shown in Fig. 4.17. In this figure, there is some noticeable contrast at the grain boundary which may be a result of the presence of accommodation strains similar to what has been observed in other studies [e.g., Valiev et al. (2000)].

Shown in Fig. 4.18 is an image of a low angle grain boundary. The FFT patterns (Fig. 4.19) show that grains A and B have the same [011] zone axis and are slightly rotated with respect to each other. As a result, the grain boundary may be described as a tilt boundary. Since the zone axes are the same, the FFT was used to determine the misorientation angle which is about 13o. This tilt boundary also displays accommodation

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Figure 4.19.

FFT of grain A and grain B from Fig. 4.18 both having a [011] zone axis and showing a slight rotation with respect to each other.

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Figure 4.20.

IFFT image of Fig. 4.18 showing the accommodation of dislocations (marked by the arrows) at the low angle grain boundary.

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strain which could have developed during the crystal joining process. In a closer examination of the boundary by producing an IFFT image (Fig. 4.20), dislocations, which are characteristic of a low angle tilt boundary, are identified (marked by the arrows). That is, when the misorientation angle is low enough between the two grains which form the tilt boundary, an array of edge dislocations can be accommodated. The dislocation positions are also consistent with the darker contrast areas at the fringe of grain B suggesting that there is an emanating strain field.

Shown in Fig. 4.21 is an example of two twin boundaries from sample no. 5 (Ni, 23 nm) within a grain having a [011] zone axis. The FFT pattern (inset) of the twins displays streaking between the spots in the [111] direction and as a result, the twinning planes are

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Figure 4.21.

HR-TEM image of two twin boundaries from sample no. 5 (Ni, 23 nm). Inset is the indexed FFT pattern and the zone axis is [011].

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Figure 4.22.

IFFT of Fig. 4.21 showing dislocations (marked by the arrows) at the twin boundary necessary to compensate for the misalignment of the lattices.

(111). In a closer examination of the IFFT image (Fig. 4.22), the twin boundary formed between A and B shows both coherent and incoherent characteristics while the twin boundary formed between B and C is generally coherent. Coherency is evident when there is alignment between the lattices or, as seen in the IFFT image (Fig. 4.22), continuity of the planes through the twin boundary. Incoherent regions in the twin boundary formed between A and B are evident due to the presence of dislocations (marked by the arrows). It is generally accepted that a dislocation array is required to compensate for the misalignment between lattices of the parent and the twin [Reed-Hill and Abbaschian (1994)].

It should be noted that not all grain boundaries were found to possess local strains which is consistent with Ebrahimi and Li (2003) who identified an unstrained low angle grain

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boundary in an electrodeposited nanocrystalline Ni-21wt.%Fe alloy. However, based on the observations made in the current study, the occurrence of strained grain boundaries is rather significant. The appearance of these local strains at grain boundaries is consistent with previous studies on nanocrystalline materials synthesized by other means, e.g., inert gas condensation followed by in-situ compacting, crystallization of amorphous precursors, mechanical alloying, and severe plastic deformation [Wunderlich et al. (1990), Ping et al.





(1995), Li et al. (2000), Valiev et al. (2000)].

The nature of these local strains at grain boundaries can be explained based on the interpretation of a non-equilibrium grain boundary that results when the conditions of compatibility based on misorientation parameters are not satisfied. Such grain boundaries have long-range elastic fields since local elastic deformation is necessary for joining of the crystals [Valiev et al. (1986)]. Notwithstanding the synthesis technique used to produce nanocrystalline materials, there is a commonality amongst the various studies to suggest that the local strains at grain boundaries are intrinsic. This implies that the local strains at grain boundaries can play a crucial role in contributing to microstrain induced XRD line broadening often seen in nanocrystalline materials, which increases as grain size is decreased [Qin and Szpunar (2005)]. As a result, it is conceivable that with decreasing grain size there is also a tendency for an increase in the occurrence of local strains at grain boundaries.

When a comparison was made between the series of images that were analyzed (e.g., Fig. 4.13-4.15), it was evident that with increasing Fe concentration in the deposit, there is

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Figure 4.23.

HR-TEM image and indexed FFT pattern (inset) of the faulted regions in the grain interior of sample no. 5 (Ni, 23 nm) having a [011] zone axis.

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Figure 4.24.

HR-TEM image and indexed FFT pattern (inset) of the faulted regions in the grain interior of sample no. 7 (Ni-16wt.%Fe, 12 nm) having a [011] zone axis.

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Figure 4.25.

HR-TEM image and indexed FFT pattern (inset) of the faulted regions in the grain interior of sample no. 9 (Ni-32wt%Fe, 10 nm) having a [011] zone axis.

also an increase in the density of lattice defects. A closer examination of these features in the grain interiors was carried out on the series of images for each of the samples. As shown in the example images (Fig. 4.23-4.25), faulted regions are evident. Fig. 4.23 displays evidence of faulted regions in sample no. 5 (Ni, 23 nm) observed with a [011] zone axis and a corresponding FFT pattern showing streaking in the [111] direction. Fig. 4.24 also shows evidence of faulted regions in sample no. 7 (Ni-16wt.%Fe, 12 nm) observed with a [011] zone axis and a corresponding FFT pattern showing streaking in the [111] direction. Sample no. 9 (Ni-32wt.%Fe, 10 nm) revealed a higher density of faulted regions, also observed with a [011] zone axis and a corresponding FFT pattern showing signs of streaking (Fig. 4.25) in the [111] direction.

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The streaking in the [111] direction, i.e., normal to the (111) planes, is characteristic for the presence of twinning [Williams and Carter (2009)]. This confirms that the defects are indeed twin stacking faults of the growth type, or “growth faults”, which are not caused by deformation processes. This is also consistent with previous studies which have noted the presence of faulted regions for both electrodeposited nanocrystalline Ni [Nakahara (1981), Yang et al. (2009), Wu et al. (2006)] and Ni-Fe alloys [Ebrahimi and Li (2003)]. In particular, it is consistent with the study of Nakahara (1981) on electrodeposited Ni who identified these faulted regions to have twinning characteristics. According to Nakahara (1981) these growth faults develop during the electrodeposition process. The normal stacking sequence on the (111) planes can change when a successive layer accidentally initiates with atoms binding to the wrong sites during the growth process. That is, if the stacking sequence for a growth fault is ABCACBA, then during the growth process the C layer atoms have mistakenly occupied the B sites resulting in the faulted sequence.

Furthermore, the addition of soluble alloying elements in fcc metals is known to decrease the stacking fault energy [Murr (1975)], including fcc Ni alloys [Nie et al. (1995)]. The addition of Fe to conventional polycrystalline Ni decreases the stacking fault energy almost linearly down to a minimum value at around 65wt.% Fe after which there is an increasing trend corresponding with a transition to the bcc structure [Charnock and Nutting (1967)]. In the present case, it is likely that there is an increased occurrence of changes in the stacking sequence on the (111) planes during the electrodeposition process for those samples with higher Fe concentration resulting in the observed increased density of these growth faults.

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4.5. X-Ray Diffraction The line positions and intensities were calculated for polycrystalline Ni with random crystallographic texture according to the Bragg law and the following X-ray diffraction intensity expression [Cullity and Stock (2001)],

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where, I is the relative integrated intensity (arbitrary units), F is the structure factor, p is the multiplicity factor, θ is the Bragg angle, and e 2 M is the temperature factor. For fcc materials F  4 f for hkl unmixed Miller indices and F  0 for hkl mixed Miller indices. All values for f and p were taken from tables presented in the book by Cullity and Stock (2001).

Table 4.3 lists the details of the calculations where f is the atomic scattering factor, λ is the wavelength of the X-ray (0.

1542 nm for Cu-K radiation), and hkl are the corresponding line positions.

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XRD patterns were generated using the method described in Chapter 3. Each pattern possesses characteristic fcc lines: (111), (200), (220), (311), (222) and (400) which are in order of increasing 2 value. Examples of XRD patterns for samples no. 5 (Ni, 23 nm) and 7 (Ni-16wt.%Fe) are shown in Fig. 4.26 and 4.27, respectively. The XRD patterns for the remaining samples can be found in Appendix C. Note that the (220) line was generally very weak for all samples except for sample no. 3 (Ni, 44 nm).

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4.5.1. Lattice Parameter The lattice parameter for each of the samples was determined using the interplanar spacing values derived from the (111) and (200) lines. The results are summarized in Table For the Ni samples, the lines remain at relatively consistent 2 positions. In the case 4.4.

of the Ni-Fe alloys, there is a general tendency for the lines to shift gradually to lower 2 positions with increasing Fe indicating that the lattice parameter is being affected by alloying. The lattice parameter values derived from the (111) lines are plotted in Fig. 4.28 as a function of Fe concentration along with those determined in other studies on materials produced using electrodeposition and conventional metallurgical processing methods.

In general, the results are similar to those obtained from studies that used electrodeposition [Jartych et al. (1992), Grimmett et al. (1993), Leith et al. (1993), Li and

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Figure 4.28.

Plot of lattice parameter vs. Fe concentration in the deposit.

Ebrahimi (2003), Fukumuro et al. (2004), and Wei (2006)]. The lattice parameter in all cases tends to increase linearly with the addition of Fe. In addition to this, the current results also agree well with metallurgically processed Ni-Fe alloys [Jette and Foote (1936), Owen et al.

(1937), Bradely et al. (1937)]. The expected linear dependence of the lattice parameter on composition is commonly known as Vegard’s law. As a solute is introduced into a solution, there is a corresponding increase or decrease in lattice parameter based on a rule of mixture,

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where, a is the lattice parameter of the solution, solvent, and solute, respectively, and x is the atom fraction of the solute. In order to calculate the lattice parameter of the solution, the lattice parameters of the solvent and solute must be given. The samples produced in this study all possess characteristics of the fcc structure. Thus, the lattice parameter for the

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solvent is taken as 0.3517 nm for fcc Ni [Jette and Foote (1936), Bradley et al. (1937), Owen et al. (1937)] and for the solute 0.3639 nm for fcc-Fe at the allotropic-transformation temperature (912oC) [Basinski et al. (1955)]. Vegard’s law values are also plotted in Fig.

4.28. In general, there is good agreement with Vegard’s law over the Fe concentration range.



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