Technique for Preparing Ultrafine ... - Wiley Online Library

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DOI: 10.1002/adma.200502619

Technique for Preparing Ultrafine Nanocrystalline Bulk Material of Pure Rare-Earth Metals** By Xiaoyan Song, Jiuxing Zhang,* Ming Yue, Erdong Li, Hong Zeng, Niandua Lu, Meiling Zhou, and Tieyong Zuo The unique configuration of unpaired 4f and 5f electrons and the rich structures of their energy levels enable rare-earth metals to possess many particular physical and chemical properties, such as high electrical conductivity, large magnetic moment, and very high complexation reactivity.[1,2] Based on these properties, the rare-earth metals and compounds have been applied extensively in permanent magnets,[3] autocatalysts,[4] superconductors,[5] etc. Demand for high-purity rareearth oxides and rare-earth metals is expected to increase particularly for use in corrosion resistance,[6] heat storage and dispersal,[7] and also in environmentally friendly applications such as in pigments for paint and plastics,[8] in cement manufacture to reduce the temperature of calcination and help save energy,[9] and in refrigeration components arising from the search for chlorofluorocarbon (CFC) replacements.[10] For the nanoscale rare-earth metals, because of the significantly increased total surface area or the grain boundary area, some new features show in the crystal structures, interface, thermodynamics, and phase transitions.[11–13] Consequently, remarkably improved optical, electronic, magnetic, and catalysis properties can be expected.[14–16] However, because of the extremely high chemical reactivity and hence the considerably rigorous equipment requirements to preserve a high purity of the product, the preparation and characterization of nanostructured pure rare-earth metals are still big challenges in nanoscience and nanotechnology. Thus, many important features of nanoscale rare-earth metals, such

– [*] Prof. J. Zhang, Prof. X. Song, Dr. M. Yue, Prof. M. Zhou, Prof. T. Zuo School of Materials Science and Engineering Key Laboratory of Advanced Functional Materials Ministry of Education, Beijing University of Technology Beijing 100022 (P.R. China) E-mail: [email protected] E. Li, H. Zeng, N. Lu School of Materials Science and Engineering Beijing University of Technology Beijing 100022 (P.R. China) [**] This work is supported by the National Natural Science Foundation of China (50401001, 50271001), the Key Project of Science & Technology Innovation Engineering, The Chinese Ministry of Education (705004), the Beijing Natural Science Foundation (2041001), and the Program of Beijing New Star of Science and Technology (2004B04). The authors thank Prof. M. Rettenmayr, from the Friedrich-Schiller-University Jena, for the very helpful discussions. Supporting Information is available online from Wiley InterScience or from the author.

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as the physical, chemical, thermal, and mechanical characteristics have rarely been reported so far. The research corresponding to these characteristics is of great importance, however, both for the development of nanoscience and nanotechnology and for extending the applications of the rareearth metals. In this consideration, we demonstrate in the present work how to prepare nanostructured bulk materials of some typical members of the rare-earth metals, laying the foundation for characterizing the physical and chemical properties of nanoscale rare-earth metals. During the past two decades, a number of techniques have been developed to synthesize nanocrystalline bulk materials, such as inert gas condensation and consolidation,[17] electrodeposition,[18] severe plastic deformation,[19] crystallization of amorphous solids,[20] surface mechanical attrition,[21] and powder metallurgy.[22–24] However, it is hard to produce nanocrystalline materials with controllable grain sizes in a wide range below 100 nm. Furthermore, in powder metallurgy for the consolidation of nanoparticles, the grain size in the synthesized bulk is generally larger than the initial particle size.[22–24] Particularly, in conventional powder metallurgy processes, a rapid coarsening of nanoparticles occurs very often, leading to the formation of grains in the submicrometer or even micrometer range. Using a new “oxygen-free” (oxygen concentration < 0.5 ppm) in-situ synthesis, where inert gas condensation was combined with spark plasma sintering (SPS) in an entirely closed system, we prepared nanocrystalline bulk material of pure rare-earth metals (Nd, Sm, Gd, and Tb) with ultrafine (< 20 nm) nanograins. Taking into account the special mechanisms of SPS consolidation, which were proposed in the literature[25,26] and were recently developed in our previous work,[27] we designed a preparation scheme with sequentially arranged processes of: amorphization of nanoparticles, nucleation and growth of the short-range ordered “clusters” inside the nanoparticle, and the complete nanocrystallization, as shown in the diagram in Figure 1. By this approach we have realized the preparation of bulk nanocrystalline materials of pure rare-earth metals. The most significant advantage of this technique is that the grain size of the resultant nanocrystalline bulk is distinctly smaller than the initial nanoparticle size, which is the first demonstration to the best of our knowledge that the traditionally accepted relationship between the size of the initial powder particles and the grain size of the sintered bulk[28] can be changed by our modified powder metallurgy technology. In virtue of this technique, nanocrystalline bulk materials with controllable grain

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Figure 1. Schematic diagram of preparing the ultrafine nanocrystalline bulk of pure rare-earth metals. a) Single-crystal nanoparticles; b) amorphization of nanoparticles in SPS consolidation process; c) nucleation and growth of the short-range ordered “clusters” inside the nanoparticle; d) complete nanocrystallization.

sizes can be obtained, which will improve the material properties in many respects; moreover, the prepared ultrafine nanocrystalline bulk materials can be used to study the special features of nanostructured materials. Here we use Sm as an example to show the preparation of the nanocrystalline bulk of pure rare-earth metals with ultrafine grain sizes by a system combining inert gas condensation with the SPS technology. In the preparation, the powder of pure Sm nanoparticles was firstly produced by inert gas condensation.[29] A transmission electron microscopy (TEM) image of the Sm nanoparticles is shown in Figure 2a. The mean diameter of the nanoparticles is measured as 38 nm. This kind of inert gas condensed Sm nanoparticles are found to have two coexisting phases, i.e., the rhombohedral and hexagonal crystal structures, respectively, as indexed from the selected area electron diffraction (SAED) pattern in the inset of Figure 2a. In the sintering process a pressure of 50 MPa, a final sintering temperature of 240 °C, and a heating rate of 50 °C min–1 were used for the Sm powder with the mean particle size of 38 nm; the optimization of these parameters was according to the displacement rate of the sintering sample. There was no holding time at the sintering temperature before the sample was cooled rapidly with cold water. A 99 % theoretical density of the sintered bulk was obtained. The amorphous structure formed in the consolidated Sm bulk, as shown in Figure 2b. It can be observed that in the consolidated amorphous bulk, the Sm nanoparticles still are separated while closely contacted. The diameter of the particles has become enlarged (ca. 90 nm on average) in the plane perpendicular to the direction of the applied pressure, indicating a compression-induced deformation. The amorphization of the Sm single crystal particles is proposed to be caused by both the SPS consolidation mechanism and the intrinsic properties of the inert gas condensed nanoparticles. As can be evaluated quantitatively,[27] when the pulsed current passes through the powder of rare-earth metals

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that have a high electrical conductivity but a relatively low heat conductivity, a high temperature can be reached rapidly in the powder particles. Since the temperatures of phase transitions of the inert gas condensed nanoparticles are reduced due to the nanometer size,[30] the high temperature induced by the Joule heating at the local site of each nanoparticle can exceed its glass-transition temperature, and can even reach the melting point at the interfaces between the nanoparticles. According to the calculations reported by Song et al.[27] and Vanmeensel et al.,[31] at local sites of the particles, the temperature induced by the Joule heating is much higher than that measured by the thermocouple placed in the wall of the graphite die. Furthermore, there is a sharp increase of the temperature from the particle body to the particle-to-particle interface. As a result, the nanoparticles are likely to have a metallic glass body and a melted boundary surface. This disordered structure becomes the “frozen-in”[32] metastable state by the rapid cooling and is thus displayed as particle-like amorphous entities. In another sintering process, we first used a pressure of 50 MPa and a heating rate of 50 °C min–1, which are the same conditions used as for the results in Figure 2b. After the temperature reached 240 °C (the densification temperature), we increased it rapidly to 400 °C at a heating rate of 300 °C min–1 and applied an in situ isothermal holding time of five minutes at 400 °C. In this way, we obtained a consolidated Sm bulk containing both amorphous and nanocrystal structures, as shown by the high-resolution TEM (HRTEM) image in Figure 2c. It can be observed that the crystallizing nucleation and the growth of the short-range ordered “clusters” took place inside the initial nanoparticle. Because there are transition regions (the short-range ordered “clusters”) between the crystallizing area and the amorphous matrix the TEM image has a low contrast. For better discrimination, we marked the corresponding zones using circles in Figure 2c. The nucleation and growth of short-range ordered clusters caused the amorphous particle to transform into a zone consisting of several domains with different crystal orientations. These domains will be the potential grains or subgrains of the nanocrystalline bulk. The SAED pattern (the inset in Fig. 2c) shows a weak inner ring overlapped with some diffraction spots, which is in accordance with the incompletely crystallized microstructure (the crystallized volume fraction was roughly calculated as 70 % from the differential scanning calorimetry (DSC) trace[33]). The indexing of the SAED pattern indicates that the nanocrystals transformed from the amorphous matrix still contain two phases, that is the rhombohedral and hexagonal crystal structures, respectively. The sizes of the nanocrystals transformed from the amorphous matrix are in a range of 5–20 nm calculated by linear intercept, implying that the ultrafine nanograined microstructures are likely to be achieved by the complete crystallization of the amorphous structures. By extending the in situ holding time at the temperature of 400 °C to 10 minutes, we obtained a nearly fully densified Sm nanocrystalline bulk, with the cor-

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COMMUNICATIONS Figure 2. a) TEM image of Sm nanoparticles prepared by inert gas condensation. Inset: the selected area electron diffraction (SAED) pattern of the nanoparticles and the indexing. b) TEM image of the amorphous structure of the Sm bulk. Inset: the enlargement of the particle interior (upper left) and the SAED pattern corresponding to the area indicated by the circle (lower right). c) High-resolution TEM (HRTEM) image of the Sm bulk containing amorphous and nanocrystalline structures. Inset: the corresponding SAED pattern and the indexing. d) TEM image of the ultrafine nanocrystalline Sm bulk. Inset: HRTEM image of the locally enlarged area, with the arrows indicating the nanograin orientations (upper left); the corresponding SAED pattern and the indexing (lower right).

responding microscopic images shown in Figure 2d. Compact ultrafine nanograins (the arrows in the upper-left inset in Fig. 2d indicate the different orientations of the nanograins) formed in the Sm bulk with a mean grain size of approximately 12 nm, and a dominant phase with the rhombohedral structure (as indexed in the lower-right inset in Fig. 2d). Applying an analogous preparation procedure as for Sm bulk, we have obtained similar nanostructures for some other members of the rare-earth metals (Nd, Gd, and Tb), confirming our proposed mechanisms illustrated in Figure 1. For comparison, we present a group of microscopic images of Gd in Figure 3, showing different nanostructure states in the preparation process. Obviously, the pure Gd nanocrystalline bulk has grain sizes much smaller than the initial nanoparticle size.

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Besides the sintering parameters, the particle size of the initial powders plays a very important role in determining the amorphous or the crystalline state of the consolidated structure and in controlling the grain size of the nanocrystalline bulk. The amorphization of nanoparticles during sintering can only take place below a critical size, which is found to be about 60 nm for Sm nanoparticles in our experiments. For the nanoparticles with a mean diameter of 75 nm, even using the same sintering conditions as those for the results in Figure 2b except that the sintering temperature was increased to 320 °C, instead of the amorphous structure, a polycrystalline bulk with a mean grain size of about 120 nm was obtained. The sintering temperature of 320 °C was determined according to the progressing of densification (characterized by the displacement rate of the sintering sample).


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COMMUNICATIONS Figure 3. a) TEM image of Gd nanoparticles prepared by inert gas condensation, with a mean diameter of about 40 nm. Inset: the corresponding SAED pattern. b) TEM image showing the amorphization of the Gd nanoparticles in the process of SPS consolidation. Inset: the corresponding SAED pattern. c) HRTEM image showing the nucleation and growth of the short-range ordered “clusters” inside the initial Gd nanoparticle. Inset: the corresponding SAED pattern. d) HRTEM image showing the complete nanocrystallization state of the Gd bulk, with the mean grain size of ca. 10 nm. Inset: the corresponding SAED pattern.

The thermodynamic and mechanical properties of the prepared ultrafine nanocrystalline Sm bulk have been characterized by the measurements of the specified heat capacity and the microhardness, respectively, as shown in Figure 4. For comparison, we also present the corresponding properties of the raw polycrystalline Sm bulk. As can be seen in Figure 4a, the specific heat capacity of the ultrafine nanocrystalline Sm bulk is remarkably increased when compared with the conventional polycrystalline material; for example, it is more than two times as high as that of the raw polycrystalline bulk in the range of the testing temperature. This reflects the substantial contribution of the highly dense nanograin boundaries to the thermodynamics of the ultrafine nanocrystalline bulk, as compared with the smaller increase of the specific heat capacity of nanocrystalline bulk materials from that of the corresponding

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polycrystalline counterparts.[34,35] The microhardness of the prepared ultrafine nanocrystalline and the raw polycrystalline Sm bulks were measured using the nanoindentation method using a load in the range of 500–4000 lN. As is known, in nanoindentation the measured value of the microhardness decreases with increasing the load in the range of lower loads and becomes constant when the load is increased to above a critical value,[36] which is 2500 lN in the case shown in Figure 4b. The comparison of the constant values for the microhardness between the ultrafine nanocrystalline and the raw polycrystalline Sm bulks is shown in the inset in Figure 4b. It can clearly be seen that the microhardness of the prepared ultrafine nanocrystalline Sm bulk is remarkably improved, i.e., it is approximately 4.5 times higher than that of the conventional polycrystalline bulk.


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Figure 4. Measurements of (a) specific heat capacities and (b) microhardness of the prepared ultrafine nanocrystalline and the raw polycrystalline Sm bulks. Inset in (b): comparison of the constant values for microhardness between the ultrafine nanocrystalline and the raw polycrystalline Sm bulks.

In summary, we have succeeded in preparing nanocrystalline bulk of pure rare-earth metals by combining inert gas condensation with SPS technology in an entirely closed system. The nanograin sizes can be tailored either by changing the nanoparticle size or by modifying the sintering parameters. Ultrafine nanocrystalline bulk can be obtained with the grain size much smaller than the initial nanoparticle size using sequentially performed preparation processes: the amorphization of nanoparticles, nucleation and growth of the short-range ordered “clusters”, and the complete nanocrystallization. This approach has the primary advantage of an “oxygen-free” in situ synthesis of the nanocrystalline bulk of pure rare-earth metals over the conventional powder metallurgy techniques and can be potentially used to prepare any nanocrystalline materials when the bulk metal precursors can be thermally evaporated. It enables comprehensive studies on a large variety of nanostructured metal materials that are highly reactive in the air.

Experimental The initial nanoparticles of the pure rare-earth metals were produced by the inert gas condensation method [29]. The production of our home-built equipment can be as high as 400 g h–1 depending on the evaporation parameters. In highly purified argon gas at a pressure of ca. 10–1 Torr, the mean size of the condensed powder particles could be controlled in the range of 20–80 nm when the applied voltage was adjusted from 10 to 20 V, the input current from 80 to 200 A, and the evaporation time from 0.5 to 3 h. In a glove box connected with the powder reservoir, the prepared powder was fed into the highstrength graphite die, which was subsequently sent to the SPS equipment (SPS-5.40-IV/ET, Sumitomo Coal Mining Co., Ltd) by a sliding rail. An external pressure was then applied to the powder through the upper and lower punches. In the sintering process, the external pressure was kept constant, and the temperature, the heating rate, and the holding time were adjusted in order to produce the consolidated bulk with different grain sizes. The density of the consolidated bulk was measured using the Archimedes method with absolute alcohol as a liquid medium. The morphology of the nanoparticles and the microstructure of the consolidated bulk of the pure rare-earth metals were observed with JEOL 2010 (operated at 200KV) and PHILIPS TEC-

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NAI F30 (operated at 300KV) transmission electron microscopes, respectively. Measurements on the special heat capacity (DSC 404C, PEJASUS) and the microhardness (Hysitron TriboIndenter) were carried out for the prepared nanocrystalline and the raw polycrystalline bulks, respectively. Received: December 7, 2005 Final version: February 13, 2006

– [1] K. A. Gschneider, in The Rare Earth (Eds: F. H. Spedding, A. H. Daane), Wiley, London 1961, Ch. 1. [2] K. A. Gschneider, L. Eyring, Handbook on the Physics and Chemistry of the Rare Earths, Vol. 1: Metals, North-Holland, Amsterdam 1978, Ch. 1–2. [3] T. Uda, K. T. Jacob, M. Hirasawa, Science 2000, 289, 2326. [4] P. G. Harrison, C. Bailey, W. Azelee, J. Catal. 1999, 186, 147. [5] C. Laurent, A. Zoltán, T. P. Gál, F. J. Braun, B. B. DiSalvo, H. Jürgen, J. Solid State Chem. 2001, 162, 90. [6] A. Amadeh, B. Pahlevani, S. Heshmati-Manesh, Corros. Sci. 2002, 44, 2321. [7] K. Yamaguchi, K. Itagaki, D.-Y. Kim, M. Ohtsuka, J. Alloys and Compounds 1995, 221, 161. [8] D. J. Fray, Science 2000, 289, 2295. [9] E. Worrell, N. Martin, L. Price, Energy 2000, 25, 1189. [10] W. Dai, B. G. Shen, D. X. Li, Z. X. Gao, J. Alloys Compd. 2000, 311, 22. [11] M. Hong, J. Kwo, A. R. Kortan, J. P. Mannaerts, A. M. Sergent, Science 1999, 283, 1897. [12] M. Yada, M. Mihara, S. Mouri, M. Kuroki, T. Kijima, Adv. Mater. 2002, 14, 309. [13] U. Sohei, S. Muhamad, I. Takayuki, N. Takashi, Synth. Met. 2005, 154, 217. [14] T. F. Rosenbaum, A. F. T. Hoekstra, Adv. Mater. 2002, 14, 247. [15] D. B. Elisabetta, T. Enrico, B. Manuela, K. Vega, K. R. Vasant, J. Eur. Ceram. Soc. 2000, 20, 2691. [16] Y. Liu, F. H. Liao, J. R. Li, S. H. Zhang, S. M. Chen, C. G. Wei, S. Gao, Scripta Mater. 2005, 54, 125. [17] H. Gleiter, Progr. Mater. Sci. 1989, 33, 223. [18] U. Erb, A. M. El-Sherik, G. Palumbo, K. T. Aust, Nanostruct. Mater. 1993, 2, 383. [19] R. Z. Valiev, A. V. Korznikor, R. R. Mulyukov, Mater. Sci. Eng. A 1993, 168, 141. [20] Y. Yoshizawa, S. Oguma, K. Yamauchi, J. Appl. Phys. 1988, 64, 6044.


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[29] N. H. Hai, R. Lemoine, S. Remboldt, M. Strand, J. E. Shield, D. Schmitter, R. H. Kraus, M. Espy, D. L. Leslie-Pelecky, J. Magn. Magn. Mater. 2005, 293, 75. [30] C. C. Yang, Q. Jiang, Acta Mater. 2005, 53, 3305. [31] K. Vanmeensel, A. Laptev, J. Hennicke, J. Vleugels, O. Biest, Acta Mater. 2005, 53, 4379. [32] K. Krištiaková, P. Švec, M. Deanko, Mater. Sci. Eng. A 2004, 375– 377, 136. [33] G. Wang, J. Shen, Q. H. Qin, J. F. Sun, Z. H. Stachurski, B. D. Zhou, J. Mater. Sci. 2005, 10, 4661. [34] K. Lu, J. T. Wang, W. D. Wei, J. Phys. D 1992, 25, 808. [35] T. Turi, U. Erb, Mater. Sci. Eng. A 1995, 204, 34. [36] R. A. Mirshams, P. Parakala, Mater. Sci. Eng. A 2004, 372, 252.

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[21] H. W. Zhang, Z. K. Hei, G. Liu, J. Lu, K. Lu, Acta Mater. 2003, 51, 1871. [22] C. Dogan, J. Rawers, D. Govier, G. Korth, Nanostruct. Mater. 1994, 4, 631. [23] T. M. Lillo, G. Korth, Nanostruct. Mater. 1998, 10, 95. [24] H. W. Zhang, R. Gopalan, T. Mukai, K. Hono, Scripta Mater. 2005, 53, 863. [25] M. Tokita, Proc. of the 2000 Powder Metallurgy World Congress (Eds: K. Kosuge, H. Nagai), Tokyo, Kyoto 2001, 729. [26] V. Mamedov, Powder Metall. 2002, 45, 322. [27] X. Y. Song, X. M. Liu, J. X. Zhang, J. Am. Ceram. Soc. 2006, 89, 494. [28] R. M. German, Powder Metallurgy & Particulate Materials Processing, Metal Powder Industries Federation, Princeton, NJ 2005, Ch. 1.

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