CHAPTER FOUR METALLIC SUBSTRATES

Transcription

1 CHAPTER FOUR METALLIC SUBSTRATES This chapter falls broadly into two parts. Firstly the preparation of textured metallic substrates is described with emphasis on the degree of crystallographic alignment. The first is Ni thin films on single crystals which can then be used as control substrates to study the deposition of further buffer layers. Also rolled Ni-based tapes have been characterised, in particular a Ni 50 Fe 50 alloy developed in Cambridge for the Brite-Euram MUST project. The chapter then focuses on the properties of such substrates in addition to the crystallographic texture, the mechanical properties, oxidation resistance, surface morphology and magnetic properties are considered. 4.1 Epitaxial Ni Thin Films Film deposition and texture Ni films, several hundred nm thick, were deposited onto MgO single crystals by DC magnetron sputtering at temperatures between C. Ar was used as the sputter gas and a pressure-distance (PD) product of around 5 Pa cm was maintained. The main parameters varied were the growth temperature and deposition rate as detailed in table 4.1. The quality of texture of the resulting film is indicated. Table 4.1 Texture of Ni films grown on {100} MgO single crystals Temp. ( C) Initial Rate (nm min -1 ) Cooling Film type (200) r.c. width A N 2 polycrystalline B N 2 epitaxial 0.18 C None epitaxial 0.31 D None epitaxial 0.72 E None epitaxial

2 Chapter Four The crystallographic alignment is assessed qualitatively by examining Ward-Wallace photographs. The single crystal MgO produces sharp spots and the Ni will give rise either to spots or lines depending on whether it is well aligned or polycrystalline. If the film is epitaxial it is also possible to confirm that it has grown with the desired {100} orientation. Figure 4.1 shows such photographs. A B {111} {200} {220} {311} {222} {400} {331} Figure 4.1 Ward-Wallace photographs of Ni films on MgO Figure 4.1a has a number of horizontal lines which correspond to reflections from planes in the Ni film. In b) however the lines are no longer apparent and a number of them have been replaced by distinct spots which have been highlighted by circles. These reflections now occur at distinct values of α, as detailed in table 4.2 below. Table 4.2 Values of 2θ θ and α for the Ward-Wallace reflections in figure 4.1 {hkl} 2θ ( ) α ( ) The values of α recorded on the photograph confirm that the Ni is oriented with {100} parallel to the sample surface and the fact that the spots are quite sharp suggests that there is good epitaxial growth. This may be further illustrated using θ-2θ goniometer scans normal to 89

3 the sample surface. Figure 4.2 shows a typical scan on an epitaxial film. The presence of just the {h00} peaks confirms that only these planes are parallel to the sample surface. 200 MgO 400 MgO Intensity (arbitrary) 200 Ni 400 Ni θ (degrees) Figure 4.2 θ-2θ scans of an epitaxial Ni film deposited on (100) MgO In order to quantify the degree of crystallographic texture of the epitaxial films, rocking curves (ω scans) were measured and are shown in figure 4.3. The FWHM values of the curves are those indicated in table 4.1. Intensity (arbitrary) Sample B Sample C Sample D Sample E ω (degrees) Figure 4.3 (200) rocking curves of Ni films deposited on MgO single crystals By examining the degree of texture of the deposited films, it is possible to assess the effects of the various deposition parameters. It is clear that the initial deposition rate and the substrate temperature are very important in the production of an epitaxial film. Sample A used a relatively high rate and low temperature and resulted in no preferential texture in the 90

4 Ni film. By decreasing the deposition rate (Samples D & E), a {100} film with FWHM of around 0.6º was produced. This degree of alignment can be improved further by increasing the deposition temperature, and a FWHM of less than 0.2º may be produced at 350ºC (B), even with a slightly higher deposition rate. These results seem to confirm that an important factor in achieving epitaxial growth is that the depositing atoms must be able to move around by surface diffusion in order to arrange themselves on the substrate template. It is not clear that cooling of the chamber walls with liquid nitrogen has a great effect as both the best and worst quality films were deposited when using it. However, it is clear that its use is not crucial in order to produce a well aligned film, and so it was not routinely used thereafter. These results confirm that the Ni has a very good {100} out-of-plane texture, but in order to be suitable for coated conductor development, it is also essential to achieve a high degree of in-plane alignment. This is investigated using X-Ray pole figures. Figure 4.4 shows the distribution of the {111} poles of Ni. The peaks are extremely sharp, the width being predominantly determined by the instrumental resolution. The position of the poles, at about 55º in ψ and 45º, 135º, 225º, 315º in φ, reveals that the film is cube textured, with no relative rotation between the unit cells of the MgO and Ni. Figure 4.4 {111} pole figure of a Ni film grown epitaxially on {100} MgO 91

5 This result is in agreement with previous work [212, 213] which indicates that Ni films grown on MgO at room temperature are polycrystalline but that the epitaxial relationship Ni{001} // MgO{001} and Ni <100> // MgO <100> can be achieved by deposition at elevated temperatures. An alternative texture (Ni (112) // MgO (010) and Ni [751] // MgO [001]) has been identified for deposition by sputtering at 400ºC [241], the pole figure shows no evidence of this up to 350ºC. Barbier et al. [215] have studied growth of the first few monolayers of Ni deposited onto MgO by evaporation. They found that there are two different epitaxial growth modes on the oxygen sublattice, an expanded Ni (001) and Ni (110) with four different in-plane orientations. The {110} orientations are responsible for relaxing the strain at the interface and are removed upon annealing at high temperatures. This supports the general observation that an elevated growth temperature is favourable in order to achieve a unique cube texture Coincidence sites The epitaxial growth mode can often be predicted or at least explained by considering the matching of the film and substrate lattices. The simplest way in which this may be done is to consider the mismatch between the unit cells. Figure 4.5 demonstrates that for Ni and MgO, the mismatch is 17%, which is very high and would generally be considered too large to facilitate simple cell matching with a strained layer. Mg MgO Ni O Ni 17% nm nm Figure 4.5 The 2-D (100) unit cells of MgO and Ni, demonstrating the large mismatch 92

6 The favourability of this orientation may be explained by considering super-cells, in which there is good matching of a smaller number of sites when several unit cells are considered as illustrated by figure 4.6 for the Ni/MgO system Figure 4.6 The MgO/Ni super-cell with matching sites at the cell corners The super-cell matching comprises 9 unit cells of Ni and 6¼ unit cells of MgO. This requires a strain of just 0.3% in order to have matching at the cell corners. Such a near coincidence site lattice model of epitaxial growth is considered in more detail in Chapter Texture Development of NiFe Tape Cold-rolled tapes The substrate mainly used in this work is a Ni 50 Fe 50 alloy. The material is commercially available in cold-rolled form in thicknesses of (nominally) 13 µm, 25 µm and 50 µm and width 10 mm or 25 mm. The tapes used in all these experiments are nominally 25 µm thick though measurements with ±1 µm accuracy have determined the thicknesses of different samples to be between µm. 93

7 Figure 4.7 shows {111} and {200} pole figures of cold-rolled NiFe. The tape strongly exhibits the copper-type texture described in section 2.3 and would therefore be expected to transform to the cube texture upon recrystallisation. a) b) Figure 4.7 a){111} and b){200} pole figures of cold rolled NiFe Texture of statically annealed tapes Annealing experiments carried out on these cold-rolled tapes in a furnace at around 500ºC, 800ºC and 1050ºC for 2 hours are reported by Glowacki et al. [216]. At the lower temperature most of the material is transformed after 2 hours, whilst at 800ºC and above the texture is fully developed. It is also worth noting that the texture is not improved by further annealing and that the same degree of alignment is achieved in all thicknesses of tape available [217]. In order to carry out more detailed recrystallisation experiments tapes have also been heated in a vacuum chamber by passing a direct current. The temperature reached is determined by the power dissipated in the tape and is explored in detail in Appendix B. Note that there is an uncertainty of around 30 C in the absolute value of the temperatures quoted, though the difference between the samples is more accurate (less than 10 C). The tape was annealed for various periods of time over a range of temperatures. The extent of development of the cube 94

8 texture was quantified by measuring {111} pole figures and comparing the intensity of the poles at ψ=55, φ=45, 135 etc. (the cube texture) with the intensity of the poles at ψ=24, φ=90, 270 (the cold rolled texture). The pole figures are shown in figure 4.8 for samples annealed for 1500 s over a range of temperatures between 500 C and 555 C and clearly demonstrate that the recrystallisation proceeds by growth of cube textured grains at the expense of those grains in the cold-rolled orientation. a) b) c) 500 C 515 C 530 C d) e) 545 C 555 C Figure 4.8 {111} NiFe pole figures showing the texture after annealing for 1500 s at various temperatures 95

9 100 % cube texture annealing time (seconds) T = 555 C T = 530 C T = 515 C T = 500 C Figure 4.9 The progressive development of texture in NiFe at various annealing temperatures Figures 4.8 and 4.9 both demonstrate a very strong dependence on temperature in the range considered. At 500 C no development of the cube texture is observed, even after annealing for over 2 hours but by increasing the temperature by just 15 C, some recrystallisation occurs within a few minutes. However, it appears that one must anneal at 530 C or above in order to fully develop the texture in a reasonable time. Figure 4.9 does not indicate a linear progression of texture development, which might be expected, but this could be due to an alternative factor. As the primary recrystallisation process occurs the defect density of the microstructure is reduced, which will probably cause a decrease in resistance and therefore a drop in temperature. This could explain why the recrystallisation process in the sample is seen to slow down. The measurements also show that the ultimate sharpness of texture achieved does not depend on the annealing temperature. A sample annealed only slightly above the recrystallisation temperature such as that shown in figure 4.8e) shows the same level of orientation as samples annealed at 670 C, 770 C and 860 C. This supports the theory that the quality of the texture that can be obtained after recrystallisation is determined at the cold rolling stage. 96

10 If a heater current of 17 A or more is used, which corresponds to a temperature of around 1100 C, the tape undergoes the process of secondary recrystallisation, outlined in section 2.3. This is characterised by the growth of large grains, which give rise to individual sharp peaks in the X-Ray rocking curve. If these specimens are then oxidised, the macroscopic grain structure can easily be observed with the naked eye. The grain size for a sample held at 1100 C for 10 min is approximately 2 mm. By analogy with the RATS technique described in section 2.4, it is likely that with further annealing the grains would grow even larger Texture of dynamically annealed tapes The NiFe tape may also be obtained in a dynamically annealed form. This material has also been characterised in detail using the standard X-Ray diffraction techniques described previously, and by EBSD. The θ-2θ scan of the NiFe tape is shown in figure As expected the (200) and (400) reflections of the cubic phase are evident. However, there are also peaks at 45.7 and 48.7 which correspond to a NiFe phase with a hexagonal structure and lattice parameters a=0.249 nm and c=0.403 nm. Whilst this phase is clearly present, the X-Ray intensities are several orders of magnitude lower than those arising from the cubic phase and so its presence should not affect the material s suitability as a substrate. 200 NiFe Intensity 1/2 (arbitrary) 400 NiFe θ (degrees) Figure 4.10 θ-2θ scan of a dynamically annealed NiFe tape. 97

11 The goniometer scan may also be used to accurately determine the lattice parameter, a, of the NiFe alloy. The best peak to use to determine a is the (400) reflection, which is resolved into Kα 1 and Kα 2 components, both of which may be used to calculate an accurate value for the plane spacing and hence the lattice parameter. Using average values of 2θ for the Kα 1 and Kα 2 peaks of ± 0.1 and ± 0.1 respectively, the resulting lattice parameter is calculated (using equation 3.2) as a=3.590 ± Å. This is higher than the value predicted by Vegard s law of 3.54 Å, but consistent with previous X-Ray measurements on Ni-Fe alloys [218]. The FWHM of the NiFe peaks was found to be very much the same as those of a single crystal silicon standard. This indicates that the size of the crystallites causes no significant broadening and that therefore the average grain size must be larger than about 0.1 µm. The crystallographic texture of the NiFe tape has been determined by studying several samples cut from different places (several metres apart) on a long reel. The tape rolling direction (RD), transverse direction (TD) and normal direction (ND) are defined to lie in the [100], [010] and [001] directions respectively, as detailed in appendix A. The distribution in the rolling direction of (001) poles about the sample normal is measured by rocking the sample about [010] this rocking curve is labelled RD. Likewise, the distribution of (001) poles in the transverse direction is labelled TD. Each sample is represented in figure 4.11 by a different colour and it can be seen that the rocking curve widths are almost exactly the same for each. Also the alignment in the rolling direction is consistently better than that in the transverse direction, and the fraction of grains misoriented by more than 7 in either direction is much smaller in the RD. 98

12 Intensity (arbitrary) RD TD ω-ω peak (degrees) Figure 4.11 (200) rocking curves of NiFe in the RD and TD from 3 different samples Figure 4.12 shows similar rocking curves but this time, the two samples are taken from completely different places on the reel. The RD curves are again sharper than the TD curves, but now there is also significant variation between the two samples. This difference in the degree of alignment along the NiFe tape reel probably arises at the tape rolling stage rather than during recrystallisation Intensity (arbitrary) TD RD ω-ω peak (degrees) Figure 4.12 (200) rocking curves of 2 NiFe tape samples. One sample is represented by the solid lines and the other by individual circles. 99

13 Again, the existence of in-plane alignment has been investigated by measuring pole figures. As the {200} pole figure 4.13(a) does not extend past 70 in ψ, it cannot be used to confirm in-plane texture. It does however demonstrate a slight difference in alignment of the (002) pole in the RD and TD. The {111} and {220} pole figures and the {220} φ scan (fig. 4.14) do demonstrate the presence of a well developed in-plane texture. a) b) c) RD Figure 4.13 Pole figures showing the distribution of a){200}, b){111} and c){220} poles in a NiFe tape Intensity (arbitrary) Phi (degrees) Figure 4.14 {220} X-Ray φ scan of a NiFe tape Table 4.3 lists the values of the half maximum widths δω RD, δω TD and δφ {220} (in degrees) for a number of different pieces of tape. Note that the values of δφ {220} have been corrected for instrumental broadening as detailed in equation

14 Table 4.3 FWHM of various X-Ray peaks for five NiFe samples. The FWHM of the φ scans is for {220} poles in the 0, 90, 180 and 270 positions Sample no ω RD ( ) ω TD ( ) ϕ 0 ( ) ϕ 180 ( ) ϕ 90 ( ) ϕ 270 ( ) average The crystallographic texture has also been studied using electron backscatter diffraction (EBSD) which measures the 3-D orientation of grains on a microscopic scale. From this data, it is possible to produce a pole figure as shown in figure An inverse pole figure (figure 4.15b) may also be constructed, showing the distribution of sample normal directions. These figures demonstrate that the cube texture is present at the surface of the sample, as well as in the bulk. Because of the relatively small sample size measured (3078 pixels, approximately 20 grains) it is difficult to compare the degree of texture to that seen in the X-Rays, but it appears similar with most pixels falling within 10 of (001). There are a small number of pixels (around 20) in apparently random orientations. These are mainly located at the grain boundaries and may arise when the electron beam falls across 2 different grains, causing the resulting Kikuchi pattern to be mis-solved. a) b) Figure 4.15 a) {100} Pole figure and b) inverse pole figure (ND) of the NiFe tape determined by EBSD measurements 101

15 The most useful representation of the EBSD data is given in figure 4.16a), the so called misorientation map. To produce this diagram the total misorientation between adjacent pixels is calculated, and the boundary is coloured depending on the level of misorientation. Ideally pixels within the same grain should be measured to have zero misorientation but, due to the finite instrumental resolution, this is not the case and the uncertainty in any measurement is approximately 1. The resulting map gives a good indication of the positions of the grain boundaries and indicates which ones are high angle type. In addition, by examining the misorientation histogram, it is possible to see which boundary angles are present in the greatest number. a) b) grain boundary frequency (%) µm angle (degrees) Figure 4.16 a) Map of the NiFe tape with boundaries coloured according to their misorientation angle. The colour coding is the same as that of the histogram, b) which shows the frequency of boundaries with a given misorientation Maps based on this boundary data may be produced to show regions of the sample which are linked through boundaries within some threshold value. Hence the threshold maps of figure 4.17 show areas of the sample which it is possible to pervade whilst crossing boundaries of no more than 2, 4, 6 and

16 a) b) c) d) µm Figure 4.17 Threshold maps showing regions of the NiFe tape connected within a) 2, b) 4, c) 6 and d) 8 The Oxford Instruments OPAL system is able to measure several adjacent regions similar to that of figure 4.17 in order to create a montage showing a larger region of the surface, approximately (0.5 mm) 2 as shown in figure This obviously produces a more representative picture, and it can be seen from the threshold maps that the same results are produced on the larger scale with the critical angle for percolation being around 5. a) b) c) d) 250 µm Figure 4.18 Threshold maps for the NiFe substrate measured over a larger area showing thresholds of a) 2, b) 4, c) 5 and d) 6. Other backscattering data measured on the TSL system in Cambridge, over the same sized area as the OPAL montage is shown in figure These results appear to show slightly better alignment, with percolation occurring at less than 4, compared with nearly 5, as measured by the OPAL system. This may be caused by real variation in the textures of the samples measured or could be associated with the resolution of the instruments used, an issue which will be addressed in more detail in section

17 a) b) c) Figure 4.19 EBSD data measured on the Cambridge TSL system showing threshold maps at a) 2, b) 4 and c) µm 4.3 Texture Development in Other Ni-based Tapes Pure Ni A 50 µm thick Ni tape from OST has also been characterised as part of this work. The θ-2θ scan is shown in figure The (200) and (400) Ni peaks are the most prominent, indicating that the cube texture is present. In addition though, the (220) reflection is also evident, indicating that the texture may not be as well developed as that of the NiFe. From the position of the (400) peak, the lattice parameter is calculated as a=3.521 ± Å. 200 Ni Intensity 1/2 (arbitrary) 220 Ni 400 Ni θ (degrees) Figure 4.20 θ-2θ scan of the Ni tape 104

18 Rocking curves about the (200) peak were carried out in both the RD and TD (see figure 4.21). As the grain size was thought to be quite large, the scans were measured using large X-Ray divergence and receiving slits (1 ) in order that a large area of the tape surface would contribute (see table 3.1). Despite this the curves are still not smooth indicating that relatively few grains contribute to the scan Intensity (arbitrary) TD RD ω-ω peak (degrees) Figure 4.21 (200) rocking curves in the rolling and transverse directions for pure Ni tape (200), (111) and (220) pole figures measured on a Ni sample are shown in figure 4.22 and confirm that the cube texture is present. The FWHM values are shown in table 4.4. a) b) c) Figure 4.22 a) {200}, b) {111} and c) {220} pole figures of Ni 105

19 Table 4.4 FWHM values for the Ni tape (in degrees) δω RD δω TD δφ (202) δφ (022) δφ (202) δφ (022) EBSD measurements have also been carried out on this material. Electron images show that the grain size is about 100 µm, consistent with the rough nature of the X-Ray rocking curves. The threshold maps for 2, 4, 6 and 8 are reproduced in figure 4.23 for a large area of the tape (1 mm x 2 mm). The threshold value is estimated to be 5.5. The misorientation histogram representing this data has a peak which lies at approximately 6.5. a) b) c) d) Figure 4.23 EBSD maps for pure Ni showing thresholds of a) 2, b) 4, c) 5 and d) NiCr A small amount of cold rolled NiCr tape (3 mm wide, 200 µm thick) was obtained and X-Ray measurements have been carried out before and after annealing. The θ-2θ plot of figure 4.24 indicates that the as-received NiCr tape is polycrystalline. From the (400), (331) and (420) peaks, the lattice parameter is estimated as ± Å. The {111} pole figure for this material was measured (figure 4.25) and shows that the material does have some preferential alignment, with a texture which resembles the cold rolling texture seen in Ni and NiFe. The out-of-plane texture seems to be quite well developed, 106

20 though further rolling is probably required in order to produce the required in-plane alignment. Intensity (arbitrary) NiCr NiCr NiCr NiCr NiCr NiCr NiCr NiCr θ (degrees) Figure 4.24 θ-2θ scan of the rolled NiCr tape Figure 4.25 {111} pole figure of cold-rolled Ni-Cr tape After annealing at 500 C and 800 C, no significant change in texture is observed which is unsurprising given that the cold rolled texture is not strongly developed. Further rolling is required before useful recrystallisation experiments can be performed. 107

21 4.4 Oxidation of The presence of a native oxide on a metallic tape will affect the crystallographic nature of the surface. For example, when depositing CeO 2 buffer layers on Ni in the presence of oxygen gas, there may be oxidation of the substrate surface, producing a thin NiO layer with a (111) texture, which can cause the buffer layer to grow with an undesirable orientation. It is also possible that a very thin oxide layer may develop on a tape at room temperature in air and so it may be necessary to ensure that such an oxide is reduced before buffer layer deposition. Simple equilibrium thermodynamics may be used in order to predict whether or not an oxide will be stable in a given atmosphere. The thermodynamics of oxide formation on the surface of a metal such as Ni is governed by the following equilibria : Ni + ½ O 2 NiO (4.1) H 2 O H 2 + ½ O 2 (4.2) H 2 + NiO Ni + H 2 O (4.3) To establish whether Ni or NiO will be stable at temperature T, consider the free energy change for 4.3. G = G 0 (H 2 O) - G 0 (NiO) + RT ln{p(h 2 O)/p(H 2 )} (4.4) where G 0 represents the standard free energy of formation of the compounds, p(h 2 O) and p(h 2 ) are the partial pressures of water and hydrogen and R is the gas constant. We can calculate the partial pressure of H 2 which, for a given background partial pressure of water vapour, will result in equilibrium. This is plotted in figure 4.26 for Ni, Fe and Pd and for comparison, a number of potential buffer layer oxides. Values of G 0 have been calculated from values in reference [219]. 108

22 log{p(h 2 )/atm} Y 2 O 3 ZrO 2 CeO 2 FeO NiO PdO Temperature (K) Figure 4.26 H 2 partial pressure required to reduce some metal oxides, assuming a background partial pressure of H 2 O of atm. Thus a partial pressure of H 2 must be supplied which is above the stability line for the metal substrate in order to reduce a surface oxide. Note that a greater H 2 partial pressure will produce a larger driving force for the reduction, which is advantageous from a kinetic viewpoint. However if an oxide buffer layer is being grown, it is important to ensure that enough oxygen is present to ensure that the buffer layer is stable, and so p(h 2 ) needs to be kept below the stability line for the buffer layer oxide. For deposition of CeO 2, the oxygen stoichiometry must be greater than CeO 1.72 in order to maintain the fluorite crystal structure. Jackson et al. [220] have calculated that at 1000 K, the fluorite phase would be unstable if p(h 2 O)/p(H 2 ) were to fall below 10-4, though at lower temperatures, this ratio is much smaller. If it is necessary to subject the substrate to a highly oxidising environment, it may still be possible to produce a textured overlayer by a process known as surface oxidation epitaxy. Whilst it is common for NiO to grow in a (111) orientation on (100) Ni at relatively low temperatures, it is possible to achieve a cube textured oxide layer at higher temperatures. This process may even be used to produce a buffer layer and is discussed in detail in Chapter

23 Even in an oxide buffered coated conductor, it is possible for oxygen to diffuse through the buffer layers during the superconductor deposition process (which usually takes place at high temperatures in an oxidising environment). Whilst the surface crystallography is no longer important at this stage heavy oxidation may cause problems if the oxide layers crack or spall from the surface. Figure 4.27 is a comparison of the oxidation behaviour of some of the Ni-based materials suitable for use as substrates [221]. The temperature is ramped up over 1 hour to 700ºC and held there whilst the percentage change in mass is measured. It can be seen that alloys of Ni with V, Fe and Cu all oxidise more rapidly than pure Ni, but that ternary alloys may be developed with more favourable oxidation properties. CuNi NiFe NiV Ni NiCrV Figure 4.27 Oxidation rate of various Ni-alloy tapes heated to 700 C. Data courtesy of A. Tuissi [221] 4.5 Mechanical Properties The mechanical properties of a material are crucial in determining its suitability for use as a coated conductor substrate, metal tapes being desirable because they are flexible and can be wound in order to produce magnets. The final conductor must be mechanically robust and as the substrate will form around 90% of the composite structure, its mechanical properties will 110

24 determine the strength of the coated conductor. Figure 4.28 shows stress-strain curves for a number of Ni-based alloys [221]. The gradient of the initial elastic region is very much the same for all the materials, indicating that they have a similar Young s Modulus. The most important parameter indicated is the yield strength of the material, the stress at which it begins to deform plastically. Pure Ni has a very low yield strength, which is to be expected given that pure materials can not exhibit strengthening mechanisms such as solid solution hardening and age hardening. The only way in which a pure metal may be strengthened is by work hardening, that is by introducing dislocations and other defects which impede dislocation motion. This is not possible for RABiTS substrates as the metal has to be annealed in order to generate the recrystallisation texture. Alloying Ni and Cu produces a material which is significantly stronger than either of the pure metals. The alloy shown, Cu 70 Ni 30, is known as Monel, and addition of small amounts of other elements may change the properties significantly. For example the alloy Monel K-500 contains small additions of Al and Ti, which can form precipitates leading to age hardening. The Ni-Fe alloy is significantly stronger than the Ni-Cu tapes, mainly due to the superior mechanical properties of pure Fe over both Ni and Cu. NiV NiFe NiCrV CuNi Ni Figure 4.28 Stress-strain curves for various Ni based alloys. Data courtesy of A. Tuissi [221] In comparison, the tensile strength of pure silver is about half that of Ni at room temperature. In addition because the melting temperature of Ag is so low (961ºC), any Ag or Ag-based alloy tape would become very soft and be liable to undergo a significant amount of creep at the elevated temperatures involved in the coated conductor production process. 111

25 4.6 Surface Morphology Roughness The optical micrographs of figure 4.29 show the surface of the NiFe tape in the as-rolled state and after annealing. After rolling, a large number of grooves are clearly present, which run in the rolling direction. After annealing they are not so pronounced but are still present. It has also been shown [217] that if the tapes are annealed in vacuum at a high temperature (1000 C) for a long time (60 hours), there is a degradation in the surface finish, through thermal etching. However, as recrystallisation experiments have shown that these conditions are not required in order to develop a suitable texture, this is not a problem. a) b) Rolling Direction 200 µm Figure 4.29 Optical micrographs showing surface of NiFe tape a) as rolled and b) after annealing Figure 4.30 shows a 2-D profile of the surface of the annealed tape, showing that the grooves have a typical trough to peak modulation of several hundred nm. This is an important factor as grapho-epitaxial growth modes can cause alternative orientations of films deposited onto rough substrates [175]. He et al. [154] found that the out-of-plane texture of buffer layers was significantly poorer when deposited on rough substrates. Therefore it is important to achieve an rms roughness of around 10 nm. This may be done by mechanical, chemical or electrochemical polishing, but alternatively it is possible to avoid such a polishing step if the work rolls are themselves very smooth [115]. 112

26 Figure 4.30 Profilometer scan demonstrating the roughness of a NiFe tape Grooving Further disruption to the surface of rolled, recrystallised tapes may be caused by grooves at grain boundaries [222]. These grooves are around 1 µm wide and up to 100 nm deep for Ni substrates annealed at 800 C, but are not so evident if the annealing is carried out at 600 C [116, 170]. As the electron micrograph of figure 2.4 shows, the groove is transferred up though the buffer layers and can cause disruption to the superconducting film. Figure 4.31 demonstrates that grooving is also a feature of the NiFe tapes, although the grooves appear to be less than 50 nm deep. Thus if they are 1 µm wide, then the angle will be smaller than 3, which is probably acceptable for subsequent film deposition. Figure 4.31 AFM image showing grain boundary grooving in the dynamically annealed NiFe substrate. 113

27 4.7 Summary The preparation of an ideal single crystal pure Ni surface to use as a control substrate has proved successful. By sputtering onto an MgO single crystal, a very sharp cube texture is obtained so long as the initial deposition rate is sufficiently low. Despite the large lattice mismatch (17%) for growth in the cube-on-cube orientation this mode is observed, probably due to the existence of a very good near coincidence site lattice match. The NiFe substrate also has a good cube texture after annealing, comparing favourably with the levels of orientation reported for other Ni-based alloys reported in the literature. The recrystallisation experiments demonstrate that the sharpness of the texture is determined during the rolling process and that the annealing temperature only determines how quickly the cube texture develops. The recrystallisation temperature for NiFe is just over 500 C and at 550 C, the texture develops fully within approximately 15 minutes. If the temperature is raised to 700 C, the process occurs within seconds. The texture of 50 µm thick pure Ni tapes is not quite so well developed, possibly due to the fact that the amount of cold work to which it has been subjected is less. This certainly seems to be the problem for the 200 µm thick NiCr alloy which shows the cold-rolled texture only very weakly and thus does not undergo primary recrystallisation. The study of other properties of various Ni-based tapes indicates that alloys often vary significantly from pure Ni. Alloys of Cu have slightly better mechanical properties, but oxidise very rapidly at 700 C. Addition of Fe enhances the mechanical properties very significantly, such that the 30 µm thick tapes used in this work are strong and easy to handle. However the drawbacks of alloying with Fe are that it worsens the oxidation resistance and also the tape remains ferromagnetic. The optimum solution seems to be to use alloys of V and Cr. V enhances the mechanical properties significantly whilst the passivating nature of Cr reduces the oxidation rate. Also both elements reduce the Curie temperature such that an alloy can be produced which is non-magnetic at 77 K. 114

28 CHAPTER FOUR...88 METALLIC SUBSTRATES EPITAXIAL NI THIN FILMS Film deposition and texture Coincidence sites TEXTURE DEVELOPMENT OF NIFE TAPE Cold-rolled tapes Texture of statically annealed tapes Texture of dynamically annealed tapes TEXTURE DEVELOPMENT IN OTHER NI-BASED TAPES Pure Ni NiCr OXIDATION OF METALLIC SUBSTRATES MECHANICAL PROPERTIES SURFACE MORPHOLOGY Roughness Grooving SUMMARY