Investigation of boundary migration during grain growth in fully ...

Investigation of boundary migration during grain growth in fully recrystallised high purity nickel Y. B. Zhang1, A. Godfrey*1, W. Liu1 and Q. Liu2 The movement of individual boundaries during grain growth after primary recrystallisation in a 96% cold rolled sample of pure nickel has been followed using electron backscatter diffraction maps of the same surface area taken after each of several annealing steps. Particular attention is given to the migration of boundaries with near S3 misorientations. During annealing, a small reduction in the length fraction the S3 boundaries was observed in the early stages of annealing. Analysis of individual boundaries shows that only those boundary segments with a misorientation angle greater than 4u move during annealing. In addition, the S3 boundaries can be divided into two groups: those with misorientations very close to the ideal twin misorientation (‘twin type’) and those with a larger deviation to the ideal twin misorientation, but still satisfying the Brandon criterion (‘neighbour type’). Of these, only neighbour type S3 boundaries are observed to move during annealing. Some implications of these observations for twin development during grain boundary engineering are discussed. Keywords: Nickel, Electron backscattering diffraction, Grain growth, CSL boundaries

Introduction There has been much interest recently in the use of nickel and nickel alloys as a substrate material for second generation high temperature superconductor (HTS) tapes, in part due to the strong cube texture that can be developed in these metals by heavy cold rolling followed by annealing treatments.1–5 During annealing, the cube texture is developed as a result of both recrystallisation and subsequent grain growth. Previous investigations, for example, have shown that, after full recrystallisation, the volume fraction of cube texture may only be y50%,6 with further high temperature annealing required to develop a very high (.95%) cube fraction. However, even after high temperature annealing, it is usual that some annealing twins are present in the microstructure, especially for nickel alloys.7,8 These twin boundaries are linked to weakening of the cube texture, and their spatial arrangement may also affect the achievable critical current density of any deposited superconducting layer.9 It is desirable therefore to avoid the presence of annealing twins in HTS substrate materials. In other applications, however, the presence of annealing twins may be desirable. For example, many studies have demonstrated that boundaries associated 1

Laboratory for Advanced Materials, Department of Materials Science and Engineering, Tsinghua University, Beijing 100084, China College of Materials Science and Engineering, Chongqing University, Chongqing 400044, China

2

*Corresponding author, email [email protected]

ß 2010 Institute of Materials, Minerals and Mining Published by Maney on behalf of the Institute Received 3 April 2009; accepted 17 April 2009 DOI 10.1179/174328409X448493

with low S value coincident site lattice (CSL) misorientations show reduced susceptibility to impurity segregation, superior ductility and higher resistance to intergranular corrosion.10–14 Based on these observations, the concept of grain boundary engineering has been proposed, in which attempts are made to control material properties by a manipulation of the grain boundary character distribution (GBCD).15,16 It is commonly accepted that, for face centred cubic metals of low to medium stacking fault energy metals, including Ni and many Ni alloys, the GBCD developed in annealed samples is often dominated by S3 boundaries, which arise due to annealing twinning.17–19 Moreover, it has been found that repetitive thermomechanical processing can be used to increase further the fraction of S3 (and S3n) boundaries.20–23 A mechanism for this was recently proposed based on the regeneration of S3 boundaries during thermomechanical processing.19,21 An important feature of this model is the presence of highly mobile non-coherent S3 boundaries (note that in this article the term ‘non-coherent’ is used to refer to any S3 boundary not on a {111} plane, rather than the term ‘incoherent twin’, which is used by some researchers to refer to a S3 boundary on a {112} plane). The evolution of twin boundaries during annealing is therefore of interest for a wide range of problems. Although it is recognised that coherent twin boundaries have a generally low mobility, very few data exist concerning the range of mobilities for non-coherent S3 boundaries. In the present experiment, therefore, the authors focus on analysing the movement of S3 boundaries during annealing and using a sequential

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Boundary migration during grain growth in high purity Ni

a 3?5 min, b 3?5z20z20 min; c 3?5z20z2z40z160 min 1 Microstructure evolution during annealing at 600uC: cube grains are shown in grey; black pixels are non-indexed points; cyan and black lines represent boundary misorientations of .2u and .15u respectively; red, magenta and green lines represent S3, S9 and S27 misorientations respectively

annealing process, investigate the mobility of noncoherent twin boundaries as a function of their misorientation characteristics.

Experimental methods The starting material was 99?996% purity Ni cold rolled to a reduction of 96% (250 mm thickness). A sequential annealing treatment was carried out at 600uC to give accumulated annealing times of 3?5, 23?5, 43?5, 83?5 and 243?5 min (steps of 3?5, 20, 20, 40 and 160 min). After each annealing step, the microstructure was investigated using a Leo 1530 scanning electron microscope equipped with a fully automated electron backscatter diffraction (EBSD) analysis system (Channel 5, HKL TechnologyOxford Instruments, UK). For this, the following procedure was used. After the first annealing step (3?5 min), the rolling plane surface was ground to SiC4000 followed by electropolishing in a 1 : 3 : 4 HClO4/ CH3COOH/C2H5OH solution at 0uC/12 V for 45 s. A microhardness indent was then made on the polished surface for use as a reference mark to facilitate reidentification of the same area after each annealing step. For all subsequent annealing steps, the sample was enclosed in a glass vacuum tube (1026 Pa) containing 0?3 atm Ar2/H2 to protect the sample surface. Orientation maps were obtained using step size of 0?8 mm over an area of 4006400 mm.2 For analysis of texture, regions were classified as being either cube or non-cube, based on a definition of 15u deviation to the ideal {001}n100.orientation. In all the EBSD orientation maps presented in this paper, cube orientations are shown in grey, and non-cube orientations are shown in white, with black pixels indicating non-indexed points. Cyan and black lines represent boundary misorientations of .2u and .15u respectively. Coincident site lattice boundaries were categorised according to the Brandon criterion (Dh,15u/S21/2),24 where Dh is calculated from the total deviation between the observed misorientation RAB and the nearest ideal CSL misorientation RS, i.e. from cos21 {K[Tr(RABRS21)21]}. The authors focus our analysis in the present paper on S3n boundaries, which are illustrated in the figures by red, magenta and green for S3, S9 and S27 respectively.

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Results Microstructure evolution Some example microstructures taken during the sequential annealing treatment are shown in Fig. 1. After annealing at 600uC for 3?5 min, the sample was fully recrystallised, with an average grain size of y10 mm and a cube texture volume fraction of 45% (Fig. 2). During

a cube volume fraction and ratio of average grain size for cube and non-cube grains; b grain growth kinetics, showing data for both surface observations from sectioned bulk material 2 Evolution of microstructure parameters during annealing

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4 Misorientation distribution for all boundaries observed to migrate, collated from all annealing steps

assumption that d&d0.27 These results suggest that the sequential annealing results are not significantly affected by any surface effects and thus confirm the validity of this technique as a means of studying of grain growth.

Evolution of boundary misorientation distribution during annealing

a all values excluding S3n misorientations; b just values for S3n (n51, 2 and 3) misorientations 3 Evolution of grain boundary misorientation distributions during annealing for total annealing times of 3?5, 43?5 and 243?5 min

subsequent annealing, some grains shrink, while some grains grow. In some places, grains of new orientation appear, either from growth below the sample surface or as a result of twinning of existing grains. After the final annealing step (243?5 min), the average grain size has coarsened to 20 mm, and the cube volume fraction increased to 63% (Fig. 2). During the entire annealing process, the average grain size for cube orientation grains is 1?5–1?7 times larger than that for the non-cube grains (Fig. 2a). This result is in agreement with previous research,25 where it was shown that cube texture fraction increases during grain growth due to advantages both of size and volume fraction. Previous in situ annealing experiments26 have demonstrated that events occurring at the surface and within the bulk of an annealed single phase aluminium alloy were similar, with similar final grain sizes in each case. To verify this for the material used in the current work, the grain growth kinetics as observed on the surface (obtained from the sequential annealing experiments) were compared with the kinetics for bulk volume (obtained from a separate series of annealed and then sectioned samples) using measurements of the grain size d from isothermal annealing data. The results are shown in Fig. 2b. For both cases (surface and bulk), the data points fall on straight lines with similar gradient. The magnitude of the gradient is approximately equal to the inverse of the grain growth exponent under the

The evolution of the grain boundary misorientation distribution as calculated from the EBSD data is shown in Fig. 3a. Note that, for clarity in this figure, data for S3, S9 and S27 boundaries are excluded due to their much higher length fractions. The increase in fraction during annealing of low angle boundaries (,15u), with a corresponding decrease in the fraction of high angle boundaries, results from the increase in cube texture fraction during grain growth. In order to analyse more carefully the evolution of the S3, S9 and S27 boundaries (S3n: n51, 2, 3), the change in length fraction for these boundaries are plotted separately (calculated as the total length of each type of S3n boundary divided by the total length of all boundaries with misorientation .2u). A slight reduction in the length density S3 boundaries is seen in the first three annealing steps, after which the value remains fairly constant. This is consistent with the results that the fraction of S3 boundaries decreases during annealing of textured copper samples.28,29

Misorientation distribution of moving boundaries From a direct comparison of the microstructure observed at the sample surface after successive annealing steps, those boundaries that move during each annealing step can be identified and their misorientation characteristics can be calculated. Collated results for all boundaries observed to move during the annealing sequence are shown as a histogram (by number, rather than length) in Fig. 4 (the grey shaded area in the figure represents values below the 2u cutoff used for definition a boundary in the EBSD data). It can be seen that no boundaries with misorientation angle ,4u are observed to move during annealing. It is important to note here that 10–15 boundaries with misorientation between 2 and 4u were seen in the examined region after each of the various annealing steps. Except for this very low angle boundary range and for twin boundaries, the shape of


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misorientation (60u n111m) compared to those of twin type S3 boundaries. Figure 5 shows distributions of the deviation angle to 60u n111m for all boundaries satisfying the Brandon criterion seen after 3?5, 43?5 and 243?5 min annealing. It can be clearly seen that most deviation angles are in the range of 0–1?5u, which suggests that most S3 boundaries are of twin type, though a certain number of neighbour type S3 boundaries are also seen.

Mobility of S3 boundaries

5 Deviation angles Dh from exact 60u n111m misorientation for all near-S3 boundaries seen in samples annealed to times of 3?5, 43?5 and 243?5 min

this distribution is similar to that for the initial boundary misorientation distribution. Some boundaries classified as S3n boundaries are, however, observed to move during annealing. These are discussed further in the section on ‘Mobility of S3 boundaries’.

Discussion Classification of S3 boundaries Annealing twins (S3 boundaries) can, in general, be classified as being of either coherent or non-coherent type. It is not possible, however, from EBSD measurements on a single surface to fully determine the twin boundary plane, although in some cases, a trace analysis can be used to rule out the possibility of a being a coherent twin. For both types of twin, the misorientation calculated as from EBSD data is expected to be very close to the theoretical value, i.e. 60u n111m, with small deviations caused only by the angular resolution of the EBSD method. In the following, therefore, both types of S3 boundary are considered as ‘twin type’ boundaries. As proposed previously,28 another type of boundary with S3 misorientation characteristics is also possible. This type of S3 boundary is referred to here as ‘neighbour type’, as such boundaries are formed by the coincidental meeting of two grains during grain growth (rather than as the direct result of a twinning process). Such neighbour type S3 boundaries normally exhibit larger deviations to the ideal S3 boundary

Although it is known that the coherent S3 annealing twin is highly immobile, much less is known about the mobility of S3 boundaries on planes other than {111}. Some indirect evidence for the mobility of non-coherent S3 boundaries can be found in the literature. For example, in one study of low temperature (540uC) annealing of copper,30 it was concluded that boundary migration of incoherent S3 boundary facets was responsible for the observed changes in the S3 boundary plane population, which took place in the absence of noticeable grain growth. In contrast, the observation of ‘island grains’ left behind during directional annealing of Ni,31 many of which have a S3 misorientation to their surrounding host grain, suggests a low mobility for these boundaries. Table 1 shows a list of all boundaries satisfying the Brandon criterion for a S3 relationship that were observed to move during at least one annealing step during the current experiment. Also shown is the total deviation angle to the exact S3 relationship. It can be seen that all the listed boundaries have deviation angles Dh larger than 3u, such that these are all of neighbour type S3 boundaries. Many of the non-moring S3 boundaries with Dh,3u were noncoherent twins (as identified from the boundary trace). The data suggest that the mobility of non-coherent twin boundaries is therefore lower than that for neighbour type S3 boundaries, even though the mobility of noncoherent twins is expected to be significantly higher than that for coherent twins. Some examples of the ways that these differences in mobility results in a change in the S3 boundary population during annealing are shown in Fig. 6. For example, in Fig. 6a and b, a neighbour type S3 boundary (just within the Brandon criterion at Dh58?6u) moves into grain A and results in the formation of another neighbour type S3 boundary (with grain C; Dh55?7u). In Fig. 6c and d, a boundary just outside the deviation angle range for a S3 boundary

Table 1 List of boundaries satisfying Brandon criterion

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Grain no. 1 orientation (Euler angles)

Grain no. 2 orientation (Euler angles)

Boundary misorientation angle

Dh (deviation to 60u n111m)

275.1 6.0 1.4 28.7 46.5 20.3 237.7 1.9 29.1 237.8 46.3 68.6 261.9 5.7 14.7 246.1 43.8 66.4 286.7 1.0 85.9 241.6 43.9 71.4 329.1 4.8 32.4 337.9 47.4 56.4 56.1 37.2 83.9 339.7 8.9 23.0

59.1 40.0 74.6 59.5 2.7 35.2 57.2 41.2 68.9 40.1 3.4 54.3 156.7 47.6 60.7 317.7 4.3 48.6 33.4 48.7 29.9 289.0 1.1 76.4 242.6 43.3 64.8 177.0 3.4 6.9 284.8 42.8 18.5 247.3 42.3 59.2

57.0 59.4 57.7 58.6 58.8 56.4 57.8 56.9 56.8 58.4 55.2 56.6

6.2 8.6 6.4 8.4 6.3 8.7 4.0 8.4 7.9 8.0 6.2 8.2


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6 Examples to show evolution of twin boundaries during grain growth: numbers in each ﬁgure indicate deviation angle Dh from ideal 60u n111m

(Dh59?5u) is seen to move during grain growth, allowing the consumption (removal) of the two twin type S3 boundaries in grain B. After grain growth, a new neighbour type S3 boundary (Dh58?0u) is formed between grains A and C. Figures 6e and f shows an example of the lengthening of a twin type S3 boundary during annealing. Here, the mobility of the medium and high angle boundary segments surrounding grain D allow the shrinking of this grain, with a consequent increase in the length of the S3 boundary between grains A and B and between A and C.

Implications of S3 mobility variations for grain boundary engineering Recently, a mechanism has been proposed for the increase in twin boundary fraction during thermomechanical processing, referred to as the ‘S3 regeneration’ model.19,21 The mechanism involves the impingement of two newly recrystallised grains that have a very similar crystal orientation and both of which have twinned during annealing. The most mobile non-coherent S3 boundaries are then expected to move and interact with existing S3 boundaries, resulting in an overall increase in the length fraction of S3 boundaries. It is clear from this description that this mechanism involves both twin type and neighbour type non-coherent S3 boundaries, and based on the present observation, the most mobile S3 boundaries should be of neighbour type. At present, the lower limit on the deviation angle from the exact S3 misorientation required to produce a highly mobile neighbour type S3 boundary is not yet established, though in the current study, no S3 boundaries with a deviation angle of less than 3?6u to the ideal 60u n111m misorientation were observed to move significantly during annealing. Under standard processing conditions (such as those used in the present study), the length fraction of neighbour type S3 boundaries is very low,

even though the sample is strongly textured. It can be suggested therefore that the success of repeated thermomechanical processing in producing an increase in the S3 boundary fraction may therefore be due in part to conditions that results in the generation of a large fraction of neighbour type S3 boundaries, with deviation angles from the exact S3 misorientation large enough to be mobile but not so large as to hinder the required S3n rearrangements.

Conclusions The migration during grain growth of individual boundary has been followed using a sequential annealing approach combined with EBSD mapping. The following conclusions can be reached from analysing the collected data: 1. The texture and microstructure evolution can be followed during grain growth by measurements on a sample surface, with grain growth kinetics similar to those seen in the interior based on a standard multisample sectioning approach. 2. S3 boundaries can be separated into twin type and neighbour type, according to the deviation angle to the ideal S3 misorientation. In the present experiment, only neighbour type twins are observed to move during grain growth, though the twin type S3 boundaries can be elongated as the result of migration of medium/high angle boundaries. 3. The results suggest that the development of neighbour type S3 boundaries may play an important role in the success of grain boundary engineering via repeated thermomechanical processing.

Acknowledgement The research was supported by the Natural Science Foundation of China (NSFC) under grant no. 50671052.


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