Download - [ACS Symposium Series] Inorganic and Organometallic Polymers II Volume 572 (Advanced Materials and Intermediates) || Synthesis of Nanocomposite Materials via Inorganic Polymer Gels

Chapter 16

Synthesis of Nanocomposite Materials via Inorganic Polymer Gels

Kenneth E. Gonsalves1, Tongsan D. Xiao2, and Gan-Moog Chow3

1Polymer Science Program, Institute of Materials Science and Department of Chemistry, and 2Connecticut Advanced Technology Center

for Precision Manufacturing, University of Connecticut, Storrs, CT 06269

3Center for Bio/Molecular Science and Engineering, Naval Research Laboratory, Washington, DC 20375

Materials with novel properties are obtained when the constituent phase morphology is reduced to the nanometer dimension. In this emerging area, investigations have involved the use of inorganic polymeric gels for the preparation of nanostructured materials containing Al-B-N or Fe-B-N. Inorganic polymer gels, or precomposite gels, were obtained by the ammonolysis of aqueous solutions of commercially available inorganic salts. The conversion of the inorganic polymer gels via thermochemical processing, resulted in the formation of nanostructured composite materials. The preparation of AlN/BN nanocomposite and FexN/BN magnetic nanocomposite materials via these chemical routes are outlined. The characterization and properties of these materials will also be discussed.

Recently, there has been considerable interest in the area of nanostructured materials. Nanostructured materials are usually referred as to materials having phases or grain structures less than 100 nm. Because of the large surface area to volume ratio of the nanostructures, a significant fraction of atoms (up to 50%) are found in the grain boundary regions. Compared to conventional materials with structures of micron size, nanostructured materials are anticipated to have superior mechanical, magnetic and physical properties [1-4]. For instance, paramagnetic and superparamagnetic nanocomposites may be produced for high density magnetic storage applications and for enhancing the efficiency of magnetic refrigeration cycles. The reduction of grain size to the nanometer regime also opens up novel processing possibilities for advanced materials. Examples are the brittle ceramics and intermetallics, which are current candidates for high temperature applications. The conventional micron size ceramics and intermetallics cannot be deformed plastically at ambient temperatures because of the small dislocation density and mobility, leading to undesirable brittle failure. However, low temperature deformability, superplasticity and ductility can be achieved for nanostructured materials that will allow an improvement of current processing methods for high temperature ceramic and intermetallic materials [4].

The starting point for the preparation of the nanostructured materials has been the synthesis of single phase metals and ceramic oxides using physical

0097-6156/94/0572-0195$08.00/0 © 1994 American Chemical Society

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In Inorganic and Organometallic Polymers II; Wisian-Neilson, P., et al.; ACS Symposium Series; American Chemical Society: Washington, DC, 1994.

196 INORGANIC AND ORGANOMETALLIC POLYMERS II

evaporation methods [1]. However, detailed studies on the synthesis of silicon based nanostructured nonoxide ceramic composite materials have been particularly intriguing [5-7]. A synergistic approach was developed, by combining interdisciplinary concepts of organometallic polymer chemistry and materials processing for the preparation of nanocomposite materials. The nanophase Si-N-C containing ceramic particles were obtained by the ultrasonic injection of a silazane precursor into (i) the beam of an industrial CO2 laser [5,6], and (ii) a hot-wall reactor [7]. The nanocomposites A1N/BN [8-9] and FexN/BN [10-11] were prepared by the thermochemical conversion of a metal-organic polymeric gel. Using organometallic precursors, multi-component metallic alloy or steels have also been produced including nanostructured Fe-Cr-Mo-C M50 steel [12], Fe-Co alloy [13], metal carbides and metal suicides [14-17]. Methodologies for these syntheses involve the designing of suitable molecular precursors that facilitate the production of nanostructured materials or composites containing phases of desired compositions and microstructures. Organometallic and inorganic precursors are particularly attractive since they provide: (i) the synthesis of materials with the ultrafine structures by assembly of small atomic or molecular clusters, (ii) phases with selected stoichiometry, and (iii) mixing of constituent phases at the molecular level. Consequently, the possibilities now exist for the synthesis of nanocomposite materials in significant quantities. These may contain a broad range of phases, such as nitrides, carbides, borides, and intermetallic phases.

In this chapter, the preparation and charaterization of A1N/BN ceramic nanocomposite and Fe xN/BN magnetic nanocomposite using inorganic polymeric gels are discussed. The latter precomposites, derived by the reaction of water soluble starting compounds, are converted to the final desired nanocomposite materials through thermochemical processing in a reactive gaseous environment The methodology presented here has the following features: (i) using an aqueous solution mixture of aluminum- and boron-containing starting compounds (ii) ammonolysis of the aqueous solution, resulting in the formation of a pre-composite gel which converts into the A1N/BN composite on further heat treatment, and (iii) a potentially economically viable synthesis, and possibly also a process of low toxicity.

1. A1N/BN Composite System

The aluminum nitride/boron nitride (A1N/BN) composite material is of considerable interest in modern technological applications because of its excellent mechanical, chemical and physical properties. Aluminum nitride and boron nitride are essentially chemically inert materials, stable at elevated temperatures. Furthermore, they are electrically nonconductive, yet possess high thermal conductivity, together with optical transparency over a wide spectral range, and have excellent dielectric properties. These properties make the A1N/BN composites suitable as an electromagnetic window material, grinding media, heat sink, as well as electric and structural materials.

The synthesis procedure is outlined in a flowchart illustrated in Figure 1. Aluminum chloride hexahydrate AlCl3*6H20, urea (NH2)2CO, and boric acid H3BO3 were dissolved in water in molecular proportions with selected stoichiometry for A l , B, and Ν as shown in Table 1.After thorough mixing, ammonia gas was bubbled into the aqueous solution mixture, while it was vigorously stirred and heated to about 90°C. After the removal of water, a precomposite gel was obtained which was then transferred into a reaction chamber for post high temperature processing to obtain the final A1N/BN nanocomposite. The precomposite material was pyrolyzed under a constant flow of ammonia at a temperature range from 500°C to 1100°C.

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16. GONSALVES ET AL. Nanocomposite Materials and Polymer Gels

Table 1; Materials Composition

197

NO. AIN Concentration BN Concentration 1 0% 100% 2 5% 95% 3 20% 80% 4 35% 65% 5 50% 50% 6 65% 35% 7 80% 20% 8 95% 5% 9 100% 0%

aluminum Chloride Dissolued in Water

Boric acid + Urea Dissolued in Water

Solution Mixed under Uigorous Stirring

Bubbling of Ammonia Gas and S loin Heating up to 90 °C

analysis of the Precomposite

Remoual of Water under Uacuum and Residue Heated up to 150 °C

Pyrolysis from 5 0 0 ° C to 1100 °C in Ammonia Gas to Obtain HIN/BN Composite

Annealing of AIN/BN Composite in Nitrogen gas up to 1 6 0 0 ° C

Figure 1. Flow chart for the preparation of AIN/BN nanocomposite materials.

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1.1 Characteristics of the Precomposite Powders

The precomposite gel has a complex polymeric gel type structure: boron bonded with urea and aluminum bonded with N-H and chlorine atoms [8]. These structures have been tentatively derived by a combination of characterization techniques including NMR, FTIR, mass-spectrometry., and x-ray diffraction studies. Thermal analysis studies performed in argon using a simultaneous high temperature TGA/DTA technique (Figure 2) show that the precomposite gel undergoes a decomposition in the 100°C to 400°C temperature range, with a total weight loss of about 80%. This weight loss is mainly due to the formation and sublimation of NH4CI, and possibly other minor species such as CO. Only one peak was evident in the DTA curve, which corresponds to the major weight loss with respect to the TGA curve. The other two smaller peaks could probably be due to die loss of excess urea from the sample and the decomposition of carbonyl groups. A combination of the TGA/DTA measurement suggested that the thermal conversion of the precomposite materials could occur at a temperature as low as 400°C to obtain die final AIN/BN composite.

1.2. Structure and morphologies of AIN/BN Composite Powders

Samples of the Al-B-N precomposite were heated in NH3 from room temperature to 500°C, 800°C, and 1100°C to obtain the AIN/BN composite. As shown in Figure 3a, the precomposite material contains various chemical bonds. The relatively sharp peaks observed included one at 1620 cm."1 corresponding to the C=0 bond, and one at 1400 cm."1 corresponding to the B-N bond [8,18-21]. The

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16. GONSALVES ET AL. Nanocomposite Materials and Polymer Gels 199

peak at 1150 cm. - 1 may arise from a C-N bond [15], while the broad band from 800 to 400 cm."1 may arise from a combination of A l bonded with N, and C species [6]. The peaks located in the 3600 to 2800 cm. ' 1 wavenumber range may be assigned to O-H and N-H bonds. The C=0 group in the precomposite gel is eliminated upon heating to 500°C in ammonia, evidence of this is provided by the M IK spectra in Figure 3b, where only three peaks are evident The peak at 1400 cm."1 is due to the B-N bond [6], while the broad peak at 3600 to 2900 cm."1 is due to a combination of O-H and N-H. The broad peak centered at 700 cm.*1 is due to Al-N bonded species. Following treatment at 800°C (Figure 3c), the small shoulder at -1650 cm."1 on the BN peak disappeared, and the N-H peak at 3400 cm."1

became much narrower and substantially decreased in relative peak height

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In thell00°C treated material (Figure 3d), two peaks at 1400 cm."1 and 700 cm."1

corresponding to the BN bond [8,18-21] and A1N bond [6] respectively, became evident There is also a small peak at 3400 cm.*1 due to the presence of a residual N-H bond in the sample. The N-H bond peak was eliminated when a sample initially heated to 1100 °C was heated further to 1400 °C in N2. Finally, the A1N peak in the FUR spectra appears to be rather broad, this is probably due to the presence of the B-N bond at a wavenumber of 790 cm."1 [18].

Figure 4 shows the XRD patterns for AIN/BN composites with varying compositions heat treated at 1100°C. When the BN concentration was below 30%, only a small amorphous BN signal was detected. For a BN content larger than 35%, broad peaks at two theta«25° and 42° were seen. These peaks were assigned to turbostratic BN, which is a two dimensional variant of hexagonal BN and has a random stacking of layers [22]. An amorphous background ranging from 25° to 45° was also found in these samples, and its intensity increased with the BN concentration. Thus, for a BN concentration larger than 35%, BN existed as a mixture of crystalline and amorphous phases. The lattice constant values of BN were not determined, but the d-spacing (002) and (101) reflections agreed with those of turbostratic BN. The A1N phase was identified to be hexagonal for concentrations of 5% to 100% A1N. When A1N concentration was 5%, hexagonal A1N peaks were very small and broad. The intensity and sharpness of A1N peaks increased with its concentration. A calculation of the lattice constant from experimental data showed a good match for A1N with the known bulk value (i.e. a=3.11À, c=4.98À, c/a=1.601). Figure 5 shows an increase of average crystalline size of AIN as a function of its concentration, with a maximum at 52 nm at 100% A1N. The crystallite size of BN was estimated to be smaller than 10 nm from the very broad (002) peaks in the XRD patterns [9].

CO c φ

s Έ v. CO

z S E 3 ζ ζ 3 3

Ami 3

-Λ .

ζ 3 t & !

J L _ A _ / v 5

100% AIN

35% AIN

100 %BN

20 30 40 50 60

2 thêta (degree)

70 80

Figure 4. XRD pattern of the composite with various AIN concentration heat treated at 1100°C.

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600

500

400 h

300 h

200 h

100 20 40 60

% AIN

80 100

Figure 5. Average crystallite size of AIN as a function of its concentration in the nanocomposite.

The structure, morphology, and crystallite size of the composite powders were correlated with compositions and processing temperatures. Figure 6 shows the bright field transmission electron microscopy (TEM) studies of the composite powders of different compositions. There was a size distribution of A1N crystallite (dark contrast), but the average crystallite size increased with increasing A1N concentration as observed in XRD measurements. Selected area diffraction of powders showed a hexagonal A1N phase in all composite samples and a turbostratic BN phase in samples where the BN concentration was larger than 35%; below this concentration, the BN phase was amorphous. The structure of the turbostratic BN phase was also confirmed from studying the lattice fringes obtained by high resolution TEM (HRTEM) studies. Although both diffraction methods did not detect a crystalline B N phase at concentrations below 35%, the existence of BN was confirmed by FITR and chemical analysis. Upon annealing, a more ordered hexagonal BN phase developed from the initial turbostratic and amorphous phases. Independent of compositions, grain growth of A1N and BN occurred, but both crystallites remained in the nanometer regime (e.g. crystallite sizes less than 60 nm for A1N, and less than 10 nm for BN) up to 1600°C, and the grain growth of BN was less sensitive to annealing temperature compared to A1N [9].

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Figure 6. Bright field TEM micrographs showing the morphological changes with the materials composition.

2. Fe xN/BN Magnetic Nanocomposite

For magnetic applications, nanocomposites are useful in high density information storage and magnetic refrigeration [2,23]. The magnetic properties of the material will change tremendously by reducing the particle size. Due to excellent magnetic properties combined with wear, oxidation and corrosion resistance over pure iron, iron nitride magnetic materials have recently received much attention [24]. The average magnetic moment per iron atom in Fe4N and Fe3N is 2.21 pu and 2.01 μη,

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respectively [24-26]. These magnetic moments of the nitrides are almost identical to the BCC α-Fe with a moment of 2.22 μη per Fe atom. These properties of iron nitrides are technologically important in potential high flux density application [24]. Magnetocaloric application requires a superparamagnetic nanocomposite whose magnetic particles are uniformly distributed in a nonmagnetic matrix phase. The superparamagnetic nanocomposite should have a large magnetic moment and a relatively small magnetocrystalline anisotropy. Iron nitride is very promising for this application. The non-magnetic matrix phase should be electrically nonconductive, but possess good thermal conductivity. Boron nitride is a good candidate because of its excellent dielectric and thermal properties [27].

The synthesis procedure for the preparation of FexN/BN is similar to that of the AIN/BN nanocomposite. Experimentally, iron chloride hexahydrate FeCl3»6H20, urea (NH2)2CO, and boric acid H3BO3 were dissolved in water in the molecular proportion 1:2:1, respectively. After thorough mixing, ammonia was bubbled into die solution with vigorous stirring until the solution was strongly basic. The reaction mixture was then heated up to 150°C, and the water was removed under vacuum. The precomposite gel was then pyrolyzed under a flow of ammonia at a temperature of about 500°C. In a variation of the above experiment, thermal chemical conversion of the precomposite gel into the final nanostructured magnetic composite was also carried out in a fluidized bed reactor.

Similar to the AIN/BN precomposite material, the as-synthesized precomposite for Fe x NBN magnetic nanocomposite also has a complex gel-type structure: boron bonded with urea, iron bonded with N-H and chlorine atoms. Also present is the NH4CI salt. Thermal analysis studies performed in argon using a simultaneous high temperature TGA/DTA technique show that the precomposite gel undergoes a decomposition in die 100 to 400 °C temperature range, with a total weight loss of about 74%. This weight loss was mainly due to the formation and sublimation of NH4CI, and possibly other minor species such as CO. Only one peak was evident in the DTA curve, which corresponds to the major weight loss with respect to the TGA measurement The other small bumps could be probably caused by the loss of excess urea from the sample and the decomposition of carbonyl groups. Combination of the TGA/DTA measurement suggested that thermal conversion of the precomposite materials could occur at a temperature as low as 400°C.

Figure 7. Bright field TEM micrographs for (a) Fe4N/BN, and (b) Fe3N/BN magnetic nanocomposite.

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Figure 7 are typical TEM micrographs of the Fe4N/BN and Fe3N/BN nanocomposite. It can be seen that the iron nitride particles (dark contrast) range from 10 nm to 200 nm in size. The size distribution was confirmed by die observation of mixture of diffused rings and spots in the electron diffraction pattern. X-ray analysis shown in Figure 8 revealed that the precomposite had a crystalline NH4CI phase. No BN and Fe x N phases were detected. When this material was heated in a furnace to 500°C in ammonia, crystalline Fe4N and a-Fe x-ray peaks were detected. The weak and broad x-ray peak located at 20=26.5° was attributed to a hexagonal BN phase. The α-Fe phase was not detected when this precomposite material was heat treated in a fluidized bed in ammonia. The x-ray analysis revealed only the existence of Fe3N, and Fe4N was not detected in this case. Thermal conversion of the precomposite in a fluidized bed reactor would provide better uniformity for the particle-gas reaction , resulting in the full conversion of a-Fe into iron nitride. The reason for the formation of Fe3N instead of Fe4N however, is not fully understood at the present time.

ο Δ Fe4N A BN • Fe3N • a-Fe ο NH4d

Precomposdte

ο

500°C in Tube Furnace

20 30 40 50 60 70

2Θ (DEGREE)

Figure 8. XRD pattern of the magnetic nanocomposite.

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Figure 9 shows the magnetization curve of the nanocrystalline Fe4N/BN and Fe3N/BN composites were made at 4.5 Κ and room temperature, respectively. The saturation magnetization and coercivity measured at 4.5 Κ were 140 emu/g and 54 Oe for Fe3N/BN, and 110 emu/g and 150 Oe for Fe4N/BN, respectively. Considering that the α-Fe, which existed in the sample as second magnetic phase, has a magnetic moment close to the average Fe moments of Fe4N and then comparing the measured value with theoretical saturation magnetization value 212 emu/g for bulk Fe4N, it was concluded that BN occupies approximately 50% of the total composition. A similar BN concentration was also obtained for Fe3N/BN composite. Although low temperature magnetic properties of these nanocrystalline Fe xN/BN are similar to their bulk materials, a rather large reduction of saturation magnetization has been observed. This is typical for fine particle magnetic materials, due to the transition for some particles from ferromagnetic to superparamagnetic state [28].

120

100

80

60

40

20

0

Magnetization curve of Fe4N/BN ι ι ι ι I ι ι ι ι I ι ι ι ι I ι ι ι ι ^ ι ι ι

<·> A A ASK

° w 300

I ι ι ι ι 1 ι ι ι ι I

10 15 20 H(kOe)

25

150 Magnetization curve of Fe3N/BN

ι ι ι ι ι ι ι ι ι ι ι

(b) I ' ' 1 1 I 1

4.5 Κ

100 h

5. 2

oo 0 oo

50 I

.1 I ι ι ι ι 1 t ι ι ι I ι ι ι ι I 10 15 20 25

H(kOe)

Figure 9. Magnetization measurement of the Fe xN/BN nanocomposite.

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ACKNOWLEDGMENT

We greatly acknowledge Professor Yide Zhang and Professor J.I. Budnick at the Physics Department of the University of Connecticut for the magnetic measurement Partial funding from Connecticut Advanced Technology Center for Precision Manufacturing is gratefully acknowledged.

R E F E R E N C E S

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3. J.C. Parker& R.W. Seigel, Nanostructured Materials, 1992, 1, 53. 4. H . Hahn & R.S. Averback, Nanostructured Materials, 1992, 1, 95. 5. K . E . Gonsalves, P.R. Strutt, T.D. Xiao, & P.G. Klemens, J. Mater. Sci., 1992,

27 (12), 3231. 6. T.D. Xiao, Ph.D. Thesis Disertation, University of Connecticut, (1991). 7. T.D. Xiao, K.E . Gonsalves, P.R. Strutt, & P.G. Klemens, J. Mater. Sci., 1993,

28, 1334. 8. T.D. Xiao, K .E . Gonsalves, & P.R. Strutt, Am. Ceram. Soc., 1993, 76(4), 987. 9. G . M . Chow, T.D. Xiao, X . Chen, & K.E. Gonsalves, J. Mater. Res., 1994,

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Oligomers & Polymers, pp. 191-97, Edited by J.F. Harrod, & R.M. Laine, Kluwer Academic Publishers, Boston, M A , 1991.

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R E C E I V E D August 1, 1994

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