J. Phys. Chem. C 2007, 111, 9437-9441
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Special Interfaces of ZrO2 Nanocrystals Sung-Wei Yeh,† Hsing-Lu Huang,‡ Dershin Gan,*,† and Pouyan Shen† Institute of Materials Science and Engineering, National Sun Yat-sen UniVersity, Kaohsiung, Taiwan, and Department of Mechanical Engineering, Chinese Military Academy, Kaohsiung, Taiwan ReceiVed: NoVember 17, 2006; In Final Form: March 30, 2007
Lattice images by transmission electron microscopy (TEM) were used to find new interfaces formed between c- and/or t-ZrO2 nanocrystals. The ZrO2 nanocrystals, prepared by ion-beam sputtering on a NaCl (100) plane, coalesced to form a straight {220}/{200} interface and a stepped {200}/{111} interface with misfit dislocations. The former is characterized by a good match of O2- lattice sites and smooth joining of low index planes across it without mismatch and dislocation. The special interfaces are applicable to fcc nanocrystals in general. The atomic structure of the new interface, however, can be studied further with an atomic-resolution TEM
1. Introduction Structure of grain boundaries is of paramount importance to mechanical and electrical properties of ceramics. Systematic studies of grain boundary in cubic (c-) zirconia have been conducted on diffusion-bonded bicrystals using high-resolution transmission electron microscopy (HRTEM) by Shibata et al.1 They focused on the structure of some ZrO2 boundaries, including a symmetric small-angle-tilt boundary and two types of symmetric-tilt Σ3 boundaries. Their HRTEM observations clarified that, when viewed along [110] zone axis, the smallangle-tilt boundary consists of a periodic array of edge dislocations of Burgers vector a/2[110] The Σ3 boundary of 70.5° tilt is atomically coherent with the {111} plane as the boundary, whereas the Σ3 boundary of 109.5° tilt is an asymmetric boundary taking alternative {111} and {115} planes as boundaries. Combining with theoretical calculations, they found that such grain boundaries possess unique coordination-deficient cation sites whose densities are correlated with the properties of high-angle grain boundaries. HRTEM was used routinely to study the crystal interfaces between metals, oxides, and semiconductors.2 The various Σ boundaries in Si and diamond were also studied.3,4 Nanocrystals were frequently observed to self-assemble epitaxially into several groups of preferred orientations on single-crystal substrates. There are many reports of nanocrystals orienting themselves on planes of ionic crystals, such as Au,5-7 Rh, Pd,8 ZnO,9 and MnO2.10. As for ZrO2 evaporated onto the NaCl (001) surface, it showed uniform diffraction rings, indicating no epitaxy relationship between the film and the NaCl substrate.11 In contrast, the radio frequency ion beam sputtering method was used successfully to prepare nanosized epitaxial c- or tetragonal (t-)ZrO2 crystallites on the NaCl (001) surface.12 The epitaxial relationships were determined to be [001]Z//[001]N, (100)Z//(11h0)N (group A), and [011]Z//[001]N, (100)Z//(100)N (group B1) or (100)Z//(010)N (group B2) between zirconia (Z) nanocrystals and NaCl (N) (001) substrate. The crystallite size, the existence of c- and/or t- phases, and the growth kinetics by rotation and coalescence of the crystallites were also reported.12 * Corresponding author. Fax: +886-7-5254099. E-mail: mail.nsysu.edu.tw. † National Sun Yat-sen University. ‡ Chinese Military Academy.
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Instead of the more difficult diffusion bonding technique,1 the self-assembly of nanocrystals on single-crystal substrates can form special boundaries automatically for an energetically favorable epitaxial state. Nanocrystals can offer special advantages in forming low-energy boundaries because they can rotate easily at moderate temperatures by Brownian motion or energy difference for their very small size, and the material and the distance that need to be transported by diffusion are also very small. Therefore, there is a good chance of finding special interface that cannot be observed in larger particles. In this work, the special interfaces of the ZrO2 nanocrystals on the NaCl (001) plane were characterized further by HRTEM, focusing on the misorientation, the coherency of the interfaces, and the local defects among the different groups of A, B1, and B2. 2. Experimental Methods Zirconia films were prepared using a radio frequency ion beam sputtering system. The target was a 3-in.-diameter Zr disk of 99.9% purity. ZrO2 nanocrystallites were deposited on the NaCl (100) cleavage plane under a working pressure of 5 × 10-4 Torr. The substrate temperature was 400 °C, the oxygen flow rate was 20 sccm, and the deposition time was 25 min. The condensed layer was stripped from the NaCl substrate into distilled water and then mounted on copper grids covered with carbon-coated collodion film for the transmission electron microscopic (TEM, JEOL 3010, under 200 and 300 kV) study. TEM lattice images were taken for 2D Fourier transformation to determine the diffraction pattern and inverse Fourier transform to characterize the details of the ZrO2 planes and interfaces. The c- and t-zirconia were indexed according to the distorted c-fluorite cell.13 The size of the crystallites thus prepared was 10-20 nm.12 3. Results The TEM electron diffraction pattern of the nano c- or t-ZrO2 crystals prepared on the NaCl (001) plane was shown and analyzed in Figure 3 of the previous report.12 The diffraction spots of group A are from nanocrystals of [001] zone axes ([001]Z//[001]N, (100)Z//(11h0)N), and those of groups B1 and B2 are two variants from nanocrystals of the same[011] zone
10.1021/jp0676437 CCC: $37.00 © 2007 American Chemical Society Published on Web 06/01/2007
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Figure 1. TEM energy-dispersive X-ray analysis result showing that no detectable amount of Cl or Na is present in the ZrO2 specimens. The Cu peaks are due to the Cu grid.
axes ([011]Z//[001]N, (100)Z//(100)N (B1) or (100)Z//(010)N (B2)), with a 90° difference in orientation between them. TEM energy-dispersive X-ray microanalysis at the interface of the nanoparticles is shown in Figure 1, which shows that there is no detectable amount of Na or Cl in ZrO2. Furthermore, in the electron diffraction pattern12 no extra diffraction spot other than those of ZrO2 was observed, indicating that there is no observable interaction between ZrO2 and the NaCl substrate. In fact, after annealing at 500 °C for 1 h there is still no observable interaction, as shown by TEM electron diffraction pattern in another report.14 3.1. A/B Interface. Figure 2 is the lattice image of two nanocrystallites in contact, and Figure 2b is the Fourier transformation from the square region in part a showing the diffraction patterns of the two nanoparticles. The diffraction pattern indexed in Figure 2c shows that the particles are of zone axes [001] (group A) and [011] (group B), as marked in the figure, and these two particles have the orientation relationship of (22h0)A//(200)B and (220)A//(022h)B, in addition to the common [001]A//[011]B zone axes. Inverse Fourier transformation was performed by selecting all of the diffraction spots in Figure 2b and indexed in Figure 2c plus the central spot. Small apertures just large enough to encircle each of them were used. The reconstructed image is shown in Figure 2d. Through this process, all of the periodic features of the original micrograph are retained but the diffuse scatterings are filtered out so that a much-improved image of lattice fringes is obtained. In Figure 2d, the special planes forming the interfaces between the two particles are marked with dashed lines. The planes in both particles are labeled, and the shared boundaries are (22h0)A/(200)B, (200)A/(111h)B, and (020)A/(1h11h)B, as indicated in the figure. There are two straight (22h0)A/(200)B interfaces and a steplike interface consisting of alternative (200)A/(111h)B and (020)A/(1h11h)B interfaces with the step angle about 100°. The (22h0)A and (200)B planes are well-aligned along the two straight (22h0)A/(200)B interfaces. Across the stepped (200)A/ (111h)B and (020)A/(1h11h)B interfaces, there are a series of misfit dislocations, as indicated Figure 2e. These dislocations are apparently to accommodate the different plane spacing between (22h0)A and (200)B planes. In Figure 2e, an extra (22h0)A plane parallel to the (22h0)A/ (200)B interface in crystallite A is marked with an arrow. This shows that the A and B nanocrystals, although very well-aligned, still have a slight misorientation between them. This is a case of imperfect-orientated attachment on specific atomic planes of nanocrystals.15 The result is similar to the imperfectly attached TiO2 nanoparticles in which a dislocation is generated in the boundary of the defect-free nanocrystals (Figure 115), as
Figure 2. (a) Lattice image of group A and B crystallites in contact. (b) Fourier transformation from the square region in part a showing the diffraction patterns. (c) The index of part b. (d) Inverse Fourier transformation from part b in which the interfaces and atomic planes of A and B crystallites are indicated. (e) Image showing the misfit dislocations. The specimen is ZrO2 deposited on NaCl (100) at 400 °C under a 20 sccm oxygen flow rate.
Special Interfaces of ZrO2 Nanocrystals
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Figure 4. Schematic drawings of the O2- and/or Zr4+ atomic planes of (a) (22h0)A and (b) (200)B with zone axes [001]A and [011]B in parallel. (c) Overlapping of parts a and b showing the good match of O2- ions.
Figure 3. (a) Lattice image of group B1 and B2 crystallites in contact. (b) The Fourier transformation from the square region in part a showing the diffraction patterns. (c) The index of part b. (d) Inverse Fourier transformation from part b in which the interfaces and atomic planes of B1 and B2 crystallites are indicated. The specimen is the same as that in Figure 2.
discussed in detail in that report. However, in our case attachment occurs on different atomic planes of nanocrystals. 3.2. B1/B2 Interface. Figure 3a is the lattice image of two nanocrystallites of groups B1 and B2 in contact and Figure 3b is the Fourier transformation from the square region in part a showing the diffraction patterns of the two nanoparticles. The diffraction pattern indexed in Figure 3c shows that both particles are of zone axis [011] and they have the orientation relationship of (02h2)B1//(2h00)B2 and (2h00)B1//(022h)B2. The image in Figure 3d, constructed from Figure 3b by inverse Fourier transform, shows the interface as indicated by the dashed line. A straight segment of the (02h2)B1/(2h00)B2 interface is observed between the two particles. Apart from the straight interface, there appears to be no other good interface so that a curved boundary is present, as indicated in the figure. 4. Discussion 4.1. {220}/{200} Interface. There are two straight (22h0)A/ (200)B interfaces in Figure 2 between crystallites of groups A
and B. The Zr4+ and O2- mixed atomic plane of (22h0)A and unmixed O2- atomic plane of (200)B are plotted in Figure 4a and b, respectively, with the directions [001]A and [011]B in parallel as zone axes. Overlapping Figure 4a and b in Figure 4c shows that the O2- ions match completely in the [110]A and [01h1]B direction. In the direction of [001]A and [011]B, there is a good match in every three O2- ions of A and every two O2ions of B, or similarly in every four O2- ions of A and every three O2- ions of B, and the mismatches are (5.7%, respectively. The match in every seven O2- ions of A and every five O2- ions of B is even less, about 1.1%, as shown in the figure. There is also a straight (02h2)B1/(2h00)B2 interface between particles of groups B1 and B2 (Figure 3). The atomic planes of the two surfaces and their mismatch are the same as that of the (22h0)A/(200)B interface in Figure 4, except that the [110]A and [01h1]B directions are the zone axes of B1 and B2 particles. The O2- ion densities of the two planes of {220} and {200} are different, as shown in Figure 4. Therefore, some local relaxation must occur and it may be accompanied by some O2vacancies to stabilize the c phase.16 The Zr4+ ions (0.084 nm), being much smaller than the O2- ions (0.138 nm) in coordination numbers of 8 and 4, respectively,17 can occupy many interstitiallike positions and minimize the energy. The local relaxation by the joining of different surfaces with underlying ionic atomic plane should be more complex, and some theoretical calculations coupled with experimental observations are necessary to see how the relaxation may be achieved. The (220)A and (022h)B planes (perpendicular to the (22h0)A/ (200)B interface) are well-aligned across the (22h0)A/(200)B interface, as shown in Figure 2d. The (020)A and (1h11h)B planes have a calculated angle difference of 9.7° due to NaCl (100) constraint, consistent with direct measurement in Figure 2. However, in Figure 2d these planes are joined smoothly, although slightly bent, across the (22h0)A/(200)B interface without forming a single dislocation. The d-spacing of the (020)A (a0/2
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Figure 5. Dimensions and angles of the (020)A and (1h11h)B planes joining coherently across the (22h0)A/(200)B interface, plotted according to Figure 2, showing that they are joined at a common length of 0.360 nm on the interface.
Figure 6. Schematic drawings of the O2- atomic planes of (a) (200)A and (b) (111h)B with zone axes [001]A and [011]B in parallel. (c) Overlapping of parts a and b showing the good match of O2- ions.
) 0.255 nm, where a0 is the lattice constant) and that of the (1h11h)B (a0 /x3 ) 0.294 nm) are different, but on the (22h0)A/ (200)B interface they are joined at the same distance of a0/x2 ) 0.360 nm, as shown in Figure 5. The angles and spacings in Figure 5 were calculated by the crystallographic relationships and are consistent with the directly measured ones. The perfect match is no accident because it is inherent in crystals of fluorite structures joined across the (22h0)A/(200)B interface. The excellent coherency may also explain its beneficial low energy state. 4.2. {200}/{111} Interface. The steplike interfaces of (200)A/ (111h)B and (020)A/(1h11h)B between particles of group A and B show some local bending of atomic planes across them. There is an angle difference of 9.7° between the (200)A and (111h)B planes imposed by the NaCl (001) substrate. Therefore, these two crystallites are to accommodate this angle difference to form
Yeh et al. a stepwise interface. There is further difficulty that the angle between (200)A and (020)A is 90° but that between (111h)B and (1h11h)B is 109.5°. To form a stepped interface using these planes, some distortions of both planes are inevitable. The angle of the steps is settled at about 100°, as in Figure 2, right in between 90° and 109.5°. The unmixed (200)A and (111h)B O2- planes are plotted in Figure 6a and b, respectively, with [001]A and [011]B pointing upward as zone axis. The overlapping figure of Figure 6c shows that in the parallel [001]A and [011]B direction the mismatch is the same as that in Figure 4c, whereas in the parallel directions of [010]A and [2h11h]B every five O2- ions of A and every two O2- ions of B can have a mismatch of about 3%. Apparently, the match is not as good as the previous {220}/{200} interface. Figure 2e shows that a series of misfit dislocations are present at the stepped interface. The (200)B plane of spacing 0.255 nm and the (22h0)A plane of spacing 0.180 nm has a large misfit. The misfit is 29.4% with respect to the (200)B plane ((0.2550.180)/0.255) and an extra (22h0)A plane is then expected every 2.4 (200)B planes (2.4x(0.255-0.180) ) 0.180). Therefore, a dislocation is expected alternatively every 2 and then 3 (200)B planes, consistent with the result in Figure 2e. There are 15 (200)B planes across the steps and 6 misfit dislocations are observed, consistent with the calculation of 15/2.4 ) 6.25 dislocations. The misfit dislocations are parallel to the [001]A and [011]B zone axis and are pure edge in nature. Being survived by the prevailed (22h0)A/(200)B boundary, these dislocations are probably immobile because of their formation via the imperfect (200)A/(111h)B attachment of the cuboctahedral nanocrystals. In fact, CeO2-ZrO2 of fluorite-type structure also tends to form cuboctahedral condensates.18,19 Besides, crystals were shown to generate dislocations by imperfectly oriented attachment of nanoparticles on a specific atomic plane.15 Such interfaceanchored dislocations are different from the mobile a/2[110] type dislocation observed in the small angle boundary in zirconia.1 Careful examination of the (220)A and (022h)B planes near the steps of the (200)A/(111h)B and (020)A/(1h11h)B interfaces reveals some local distortions due to the misfit dislocations and the 100° compromise as mentioned. In contrast, these two planes are very well-aligned and straight without any dislocation across the more-stable (22h0)A/(200)B tilt boundary in Figure 2e. The steplike interface of (200)A/(111h)B and (020)A/(1h11h)B is clearly not as stable as the (22h0)A/(200)B interface. First, the oxygen ions do not match as well. Second, the angle of the steps of 100° formed by bending {200}A planes and {111}B planes induces some local distortion. Third, the misfit dislocations and the distortion of the (220)A and (022h)B planes near the steps would still raise the interfacial energy. However, the stepwise {200}/{111} interface is still a low-energy interface because it was observed frequently in this research for the attached nanocrystals. 4.3. Concluding Remarks. Theoretically, low index planes in [001]A zone axes of A particle and low index planes in [011]B zone axes of B particle can meet to form interfaces, but only the two types of interfaces ({220}/{200}, {200}/{111}) were observed. It must be emphasized that these particles are fastened in orientation by the underlying NaCl (001) plane and cannot rotate. The observed special interfaces are what these particles can best manage under the circumstances. Unlike the artificial tilt Σ boundaries of bulk ZrO2,1 the present boundaries of nanocrystals are formed by different zirconia atomic planes for surprisingly good matches under the constraint of NaCl. With
Special Interfaces of ZrO2 Nanocrystals the help of NaCl (001) plane substrate, many such interfaces were formed simultaneously and were therefore recognized in this experiment. Because the diffraction pattern of c- or t-ZrO2 is similar to the fcc crystals, the interfaces should be applicable to nanocrystals of fcc structure. In crystals larger than these nanocrystals, the formation of these interfaces may require some degree of rotation and some diffusional transfer of material that must be difficult due to their larger size. However, although the new interfaces are identified in this report, atomic-resolution TEM and image simulation are necessary to study the atomic structure of the interfaces.1-4 5. Conclusions 1. The nanoparticles in group A and B were impinged and coalesced to form {220}/{200} and {200}/{111} boundaries; whereas the particles in group B1 and B2 form {220}/{200} interface. 2. The {220}/{200} interface of zirconia is found to be of especially low interfacial energy. The O2- lattice sites are shown to have good matches, the {220} planes across the interface are perfectly aligned, and the {200} and {220} planes across the interface are joined smoothly without mismatch strain and dislocations. 3. In the {200}/{111} interface of zirconia, the O2- lattices are not matched as well. These two planes form a stepped interface with distorted interplane angles. Misfit dislocations are also founded at the steps.
J. Phys. Chem. C, Vol. 111, No. 26, 2007 9441 Acknowledgment. This research was supported by National Science Council, Taiwan, ROC under contract NSC93-2120M-110-001, and partly by Center for Nanoscience and Nanotechnology at NSYSU. References and Notes (1) Shibata, N.; Yamamoto, T.; Ikuhara, Y.; Sakuma, T. J. Electron Microsc. 2001, 50, 429. (2) Ichinose, H. Sci. Technol. AdV. Mater. 2001, 1, 11. (3) Sawada, H.; Ichinose, H. Scr. Mater. 2001, 44, 2327. (4) Sawada, H.; Ichinose, H.; Kohyama, M. Scr. Mater. 2004, 51, 689. (5) Masson, A.; Me`tois, J. J.; Kern, R. Surf. Sci. 1971, 27, 463. (6) Kuo, L. Y.; Shen, P. Surf. Sci. 1997, 373, L350. (7) Chang, E. C.; Lu, F. H.; Shieu, F. S. Mater. Chem. Phys. 2001, 70, 137. (8) Reniers, F.; Delplancke, M. P.; Asskali, A.; Rooryck, V.; Van Sinay, O. Appl. Surf. Sci. 1996, 92, 35. (9) Henley, S. J.; Ashfold, M. N. R.; Cherns, D. Surf. Coat. Technol. 2004, 177-178, 271. (10) Nilsen, O.; Foss, S.; Fjellvag, H.; Kiekshus, A. Thin Solid Films 2004, 468, 65. (11) El-Shanshoury, I. A.; Rudenko, V. A.; Ibrahim, I. A. J. Am. Ceram. Soc. 1970, 53 (5), 264. (12) Yeh, S. W.; Huang, H. L.; Gan, D.; Shen, P. J. Cryst. Growth 2006, 289, 690. (13) Teufer, G. Acta Crystallogr. 1962, 15, 1187. (14) Yeh, S. W.; Hsieh, T. Y.; Huang, H. L.; Gan, D.; Shen, P. Mater. Sci. Eng. 2007, A452-453, 313 (15) Penn, R. L.; Banfield, J. F. Science 1998, 281, 969. (16) Fabris, S.; Paxton, A. T.; Finnis, M. W. Acta. Mater. 2002, 50, 5171. (17) Shannon, R. D. Acta. Crystallogr. 1976, A32, 751. (18) Lee, W. H.; Shen, P. J. Cryst. Growth 1999, 205, 169. (19) Kuo, L. Y.; Shen, P. Mater. Sci. Eng. 2000, A277, 258.
