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Mechanism of reactive sintering of MgAlB14 by pulse electric current
Int. Journal of Refractory Metals & Hard Materials 27 (2009) 556–563
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Int. Journal of Refractory Metals & Hard Materials journal homepage: www.elsevier.com/locate/IJRMHM
Mechanism of reactive sintering of MgAlB14 by pulse electric current David J. Roberts 1, Jinfeng Zhao, Zuhair A. Munir * Department of Chemical Engineering and Materials Science, University of California, Davis, CA 95616, USA
a r t i c l e
i n f o
Article history: Received 23 June 2008 Accepted 11 July 2008
Keywords: Aluminum magnesium boride AlMgB14 Ultrahard boride Reactive sintering
a b s t r a c t The reactive sintering of the hard boride AlMgB14 from the elements was investigated using the pulse electric current sintering (PECS) method. The mechanism of formation of AlMgB14 was investigated over the temperature range 600–1400 °C. The formation of AlNgB14 begins at as low a temperature as 800 °C and is essentially complete at 1325 °C. Thus the prospect of forming nanostructured AlNgB14 is not high since earlier nucleated grains continue to grow up to the maximum sintering temperature. Under optimum conditions, the products contained AlMgB14 with a 93 vol%, the remaining phases include spinel and impurity from milling. The product had a hardness of 26.1 GPa and a fracture toughness of 3.1 MPa m1/2. Ó 2008 Elsevier Ltd. All rights reserved.
1. Introduction The search for new hard materials for use in cutting tool, heat sink, abrasive, and coating applications has in the past few decades focused on compounds based on the light elements B, C, and N. Ternary compounds of these elements with a structure similar to diamond are expected to have properties similar to diamond and cubic boron nitride [1,2]. The discovery by Cook et al. [3] of an exceptionally high hardness in AlMgB14 has generated renewed interest in boron-rich compounds as hard materials. These authors reported a hardness of as high as 35 GPa for this boride; with the addition of a second phase, the hardness increases further and approaches the value for cubic BN. Such high hardness, however, is not predicted from theoretical calculations. Lowther [4], using ab initio calculations, showed that the bulk modulus of AlMgB14 is lower than that of B4C, a compound also based on the boron icosahedra structure [5,6], but with a considerably lower hardness (<20 GPa [7]). Using the concept of a ‘‘chemical hardness” [8], the author attributed the hardness difference between AlMgB14 and B4C to the difference in the location of the Fermi level relative to the energy gap [4]. Peters et al. attributed the high hardness of AlMgB14 composites to the intrinsic matrix property and the microstructure [9]. Related to the role of microstructure, observed variation in hardness values were attributed to differences in the powder processing of this boride [10]. The effect of sample preparation on the property of the boride has been recently demonstrated [11]. A significant difference in the bulk modulus
between samples prepared by two methods was attributed to possible differences in impurity incorporation, an observation that underscores the importance of the processing step. The boride AlMgB14 has been prepared by several methods. Okada et al. [12] and Higashi et al. [13] prepared single crystals of AlMgB14 from metal salts and characterized the crystalline structure of the boride and also measured its Vickers hardness, reporting values for the latter in the range 20.2–28.6 GPa. The boride has also been prepared by high-energy milling of the elements followed by hot-pressing or sintering [11]. Takeda et al. [14] used metal hexaborides as precursors for the preparation of AlMgB14 employing the pulsed electric current sintering method (PECS). In addition, thin films of this boride have been synthesized by pulse laser deposition (PLD) [15]. Although the emphasis of the majority of recent investigations on AlMgB14 is on its mechanical properties [10,15,16–18], other investigations have focused on the electrical and thermoelectrical (TE) properties of this boride. Takeda et al. [14] concluded that the boride has little prospect as a TE material because of low electrical conductivity, despite a relatively high Seebeck coefficient. In this paper we present results from an investigation on the simultaneous formation and consolidation of AlMgB14 from the element by the pulse electric current sintering (PECS) method. Takeda et al. [14] used this method, but their starting materials are different and their aim was the investigation of the thermoelectric properties of AlMgB14. 2. Experimental materials and methods
* Corresponding author. E-mail address: [email protected] (Z.A. Munir). 1 Present address: Lawrence Livermore National Laboratory, 7000 East Avenue, Livermore, CA 94551, USA. 0263-4368/$ - see front matter Ó 2008 Elsevier Ltd. All rights reserved. doi:10.1016/j.ijrmhm.2008.07.009
The starting materials were elemental aluminum, magnesium, and boron. The aluminum powder was 99.97% pure (metals basis), 325 mesh (maximum particle size approximately 44 lm)
D.J. Roberts et al. / Int. Journal of Refractory Metals & Hard Materials 27 (2009) 556–563
obtained from Alfa Aesar (Ward Hill, MA). The magnesium powder was 99.8% pure, 325 mesh, and also obtained from Alfa Aesar. Three types of starting boron were used: amorphous boron powder (99% pure with a sieve classification of 325 mesh), crystalline boron powder (90% pure with an average particle size of <5 lm), and course powder from pulverized boron pieces (<20 mm). All boron materials were obtained from Alfa Aesar. The latter type was used on the premise that the large pieces have smaller amounts of surface oxide contamination. Because oxygen contamination of the preliminary powders was the main concern in producing a single-phase product (i.e., void of the spinel phase MgAl2O4), the storing and handling of powders were done in a glove box under an argon atmosphere. The oxygen level in the glove box was 60.20 ppm. The only case in which the powders were not handled under a controlled atmosphere was during the X-ray diffraction analysis of the starting Al, Mg and B powders and the milled, pre-sintered powders. These powders were initially scanned while sealed in plastic wrap. However, the plastic wrap caused too much interference to give clear XRD patterns of the milled powders. So the plastic was removed and the exposed powders were scanned again in the X-ray diffractometer. From preliminary analysis, it was determined that the source of the oxygen contamination was most likely to be the boron powder. Hence, several different boron-processing methods were tested in an attempt to find the method that yielded a final boride product with the minimum amount of spinel contamination. The aluminum and magnesium powders were used directly from the manufactures containers. The processing methods hereafter will be referred to by the boron type, as can be seen in Table 1. Stoichiometric amounts of various boron powders, aluminum, and magnesium powders (total mass, 7 g) were subjected to high-energy milling in a planetary mill (Fritsch Pulversette 5). The jars were Sialon-lined and the balls were zirconia (10 mm diameter). A charge ratio (ball to powder mass ratio) of 14 was used. The powders were introduced into the jars inside the glove box. The jars were then sealed with parafilm and electrical tape and removed from the glove box. The sealed jars were placed in the planetary mill, which was then operated at a rotation speed of 250 rpm with a milling cycle of 1 h milling, 1 h rest for a total of 20 cycles (40 h total time; 20 h milling time). After the milling was completed, the powders were removed from the jars inside the glove box. In nearly all cases, milling was done in dry condition. However, in a few cases the powder mixture was milled in methanol. For these, the milling time was 8 h. In cases of wet milling, the powders were subsequently dried under vacuum for 72 h. As indicated in Table 1, boron powders (B-3) were annealed in vacuum to reduce the oxygen contamination. In the presence of moisture in the air, boron oxide transforms to the hydroxide
B2 O3 ðsÞ þ 3H2 OðgÞ ¼ 2BðOHÞ3 ðsÞ
having a DG < 0 at 300 °C. The amorphous boron powder (99% pure) was annealed in vacuum inside an alumina crucible that had been sprayed with boron nitride. Annealing was done in a Brew vacuum furnace for 2 h at 1500 °C under an ambient pressure of approximately 6.7 103 Pa. At this temperature, the vapor pressure of B2O3 is 1.88 103 atm. The annealing treatment also resulted in the crystallization of the amorphous boron. Using a similar treatment on amorphous boron with a purity of 99.7 wt%, Brandstötter and Lengauer [19] reported a decrease on oxygen content from 10,000 to 10 ppm after vacuum treatment at the same conditions, 1500 °C for 2 h. Also, as shown in Table 1, the amorphous boron (B-1) was milled in methanol. This was attempted on the basis of an unpublished research indicating that B2O3 reacts with methanol to produce boron esters which volatilize during the subsequent drying of milled powder [20]. The microstructure of the milled powder mixture is shown in Fig. 1. The figure shows agglomerated particles some as large as about 4 lm. X-ray analysis of the milled samples, Fig. 2, showed the presence of the reactant elements only with no indication of product formation during the milling. The figure shows peaks belonging to boron, indicating that milling had resulted in the crystallization of the initial amorphous phase. Grain size analysis on the milled powders was made by the Williamson-Hall method [21]. The result gave grain sizes for Al and Mg of 420 and 46 nm, respectively. In view of the limitation of this method with respect to grain size (100 nm), the value obtained for Al is taken as only approximate. The limited number of boron peaks and their low intensity made the grain size evaluation of this element by this method uncertain. Approximately 2.5 g of the milled powder mixtures were placed in a graphite die which was then placed in a zip-sealed plastic bag and transferred from the glove box to the spark plasma sintering apparatus. The die and plungers used for sintering the milled powders were made of high-density graphite. The die was 3.81 cm high with an outer diameter of 4.45 cm and an inner diameter of 1.91 cm. The plungers were 2.54 cm high and 1.91 cm in diameter. The inside of the die was bored an additional (approximately) 250 lm to accommodate a liner. Graphite foil approximately 125 lm thick was used as a liner inside the die. The liner acted as a barrier to inhibit the sintering powders from bonding to the graphite die. To keep the powders from bonding to the plungers, graphite foil was cut into circles about 1.91 cm in diameter and a double layer was placed between the powders and the plungers.
ð1Þ
Reaction (1) has a DG = 60.9 kJ at 25 °C. And the formed hydroxide is relatively volatile, with the reaction
BðOHÞ3 ðsÞ ¼ BðOHÞ3 ðgÞ
ð2Þ
Table 1 Boron samples and their treatment Boron designation
Source
Treatment
B-1 B-2 B-3
99.0%; amorphous 90%; crystalline 99.0%; amorphous
B-4 B-5
99.0%; amorphous 99.5%; crystalline; crushed pieces
None None Vacuum annealed (1500 °C) Milled in methanol; 8 h None
557
Fig. 1. SEM image of milled reactants Al, Mg, and B.
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The simultaneous synthesis and consolidation of AlMgB14 from powder mixtures was performed in a spark plasma sintering apparatus (SPS). The apparatus is the SPS-1050 made by the Sumitomo Mining Company, Ltd. (Kawasaki, Japan). The power supply has a 15 V, 5000 A capacity. The current generated is a pulsed DC current with a cycle of 12 pulses on and 2 pulses off, with each pulse having a duration of about 3 ms. The uniaxial press can provide up to 100 kN of force. The die containing the powder mixture was taken out of the zip bag and placed in the SPS, which was quickly closed and evacuated. Before sintering, the powders in the die were cold-pressed for approximately 20 s under a load of 10 kN (35 MPa). The load was then reduced to 5 kN (17.5 MPa). Once the SPS chamber pressure reached 10 Pa, sintering was initiated. A current was then provided to increase the temperature in the following manner: from room temperature to 600 °C in 2 min at a pressure of 17.5 MPa, held at 600 °C for 3 min as the pressure was increased to 63 MPa, then from 600 to 1500 °C the temperature was increased at a rate of 100 °C min1. The samples were than held at the final temperature for 5 min, after which the power was turned off and the samples were allowed to cool down. Subsequent experiments were carried out under a higher pressure of 123 MPa. A pyrometer (CHINO, model IR-AH, Japan), focused on the outer wall of the die, was used to measure the temperature of the system. The density of the samples was calculated from geometric and mass determinations and also by the Archimedes method. The microstructure of the samples was determined by optical and scanning electron microscopy. An FEI XL30-SFEG scanning electron microscope (SEM) with energy dispersive X-ray spectroscopy (EDX) capabilities was used to observe the polished surface of the mounted samples. A portion of the unmounted sample was
Fig. 3. SEM image of sample made with annealed boron and sintered at 1500 °C.
broken off and the fractured surface of the sintered samples was also observed under the SEM. Chemical analysis of the surface was done using the EDX function of the SEM. For X-ray analysis a portion of the sample was pulverized. Analysis was made with a Scintag X-ray Diffractometer (XRD), model XDS 2000, with Cu Ka radiation, k = 1.540562 nm. The scanning step size was 0.02 degrees for the quick scan with a hold time of 1.5 s per step, and 0.005 degrees with a hold time of 4.5 s per step for the slow scan. Data from the XRD long scan times were used to
Al Mg
Relative Intensity
β-B B
10
20
30
40
50
60
70
2θ (degrees) Fig. 2. X-ray analysis of milled reactants.
Table 2 Phase analyses (vol%) of samples sintered at 1500 °C with different boron types Sample B-1, B-2, B-3, B-5, B-4, B-1,
Al, Al, Al, Al, Al, Al,
Mg Mg Mg Mg Mg Mg (wet-milled)
AlMgB14
Spinel
Al
Mg
B
AlMgB4
Mg3(BO3)2
ZrB2
Si3N4
85.6 22.4 92.6 29.5 85.1 9.2
6.8 56.8 3.3 22.0 6.3 78.6
1.3 1.6 1.2 3.3 0.7 5.0
0.1 0 0.2 1.0 0 0
3.0 11.6 1.1 8.4 3.5 4.8
0.2 0 0.8 1.0 0 0.4
1.6 4.0 0.4 12.2 2.4 0.2
0.3 0.3 0.2 6.1 1.4 0.4
1.1 3.4 0.4 16.5 0.7 1.5
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Fig. 4. EDX analyses of regions in Fig. 3: (a) The main, dark, phase; (b) gray region; and (c) bright, nearly white region.
Table 4 Grain size, density, and hardness of AlMgB14 samples sintered at various temperatures
Table 3 EDX analyses (at.%) of different regions in Fig. 3
Dark region (Fig. 4a) Off white (Fig. 4b) Bright white (Fig. 4c)
B
O
Mg
Al
Zr
Temperature (°C)
Mean grain Size (lm)
Density (g cm3)
Hardness (GPa)
92.7 0 92.79
0 39.15 0
4.22 22.53 3.29
3.04 38.31 2.50
0 0 1.42
1500 1400 1325 1250 1000 800 600
4.76 3.87 3.44 1.75 1.58 0.37 0.36
2.66 2.68 2.75 2.32 – 1.56 1.39
23.9 26.1 24.6
perform a grain size analysis on the pulverized samples using the Williamson-Hall method [21].
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The Vickers hardness (Hv) was determined, using a 1 kg load and a 15 s dwell time on a Buehler Micromet 2004 microhardness tester. Indents were made in 10 random locations on the sample. The sample fracture toughness was calculated from the length of the cracks formed at the corners of the indents [22]. 3. Results and discussion The initial aim was to determine which of the types of boron provides the best result. Regardless of the type of boron used or the method of milling, all samples sintered at 1500 °C (at 63 MPa) contained the phase AlMgB14, as can be seen from Table 2 which lists representative results for the different conditions. The phase compositions in the table are in vol%, calculated from X-ray data using the PowderCell software. The results provide a wide range of AlMgB14 content, from about 10 vol% to a high of about 93 vol%. They indicate that wet milling of all reactants (in methanol) results in samples with the least amount of the desired boride, probably because of ineffective mixing and folding of the reactant metals and the possible oxygen contamination from water in the alcohol. The sample that gave the highest boride content and least amount of spinel was that in which annealed boron (B-3) was used. SEM image of a sample made with annealed boron and reactively sintered at 1500 °C under a pressure of 63 MPa is shown in Fig. 3. It should be recalled this sample contained the highest vol% of the desired boride, 93%. EDX analysis on this sample showed the main phase (dark region) contains Al, Mg, and B, Fig. 4a, and thus represents the desired boride phase. The inclusions in Fig. 3 are represented by two regions: a relatively major one with an off-white appearance and a minor region with bright white appearance. EDX analysis of the former, Fig. 4b, indicates that this is the spinel phase, while the latter, Fig. 4c, is a zirconium-rich phase, possibly a boride originating from contamination from the milling balls. Because of the small size of these latter contaminants, the EDX analysis includes peaks of the background, the boride phase. The results of the EDX analysis are presented in Table 3. As stated above, the vol% of AlMgB14 in this sample was about 93%, which would mean that the other phases (including the spinel) constitute about 20% on the basis of area, as depicted in the SEM image (Fig. 3). The average density of the samples providing the best results from Table 4 was determined as 2.66 and 2.62 by the geometric and Archimedes methods, respectively. The theoretical density of AlMgB14 is reported as 2.6 g cm3. To understand the process of formation and densification of AlMgB14 from the elements we conducted SPS experiments to intermediate temperatures. Typical SPS temperature and displacement profiles for a sample using annealed boron are shown in Fig. 5. The displacement curve reflects uniaxial changes in the sample dimension and thus relate to shrinkage; the displacement is positive when the sample shrinks and negative when it expands. The lowest temperature measured by the pyrometer is 600 °C and thus the presence of the flat portion on the initial temperature profile, ending at point 1a in Fig. 5. The increase in shrinkage at point 1b is the result of increasing the applied pressure from 17.5 to 63 MPa. On the other hand, the considerable shrinkage beginning at point 2b (corresponding to 1250 °C) is due to the simultaneous synthesis and densification of the boride. The temperatures indicated along the displacement curve are those used to halt the process for investigation of the synthesis and consolidation process. Samples were heated along the temperature curve of Fig. 5 and then the heating is stopped at selected temperatures of 600, 800, 1000, 1250, 1325 and 1400 °C. For the case of 1325 and 1400 °C, the pressure utilized was 123 MPa. X-ray analyses were then performed on the samples to determine phase composition. Fig. 6a
shows the results for 600 and 800 °C, Fig. 6b shows the results for 1000 and 1250 °C, and Fig. 6c shows the results for 1325 and 1400 °C. As can be seen from Fig. 6a, when the milled powders are heated in the SPS to 600 °C, partial reaction of the elements takes place, producing the lower boride phase, AlMgB4. No evidence of the AlMgB14 can be seen at this temperature. At this temperature, the reaction to form any boride is not complete, as seen by the presence of peaks for Al and Mg. However, when the sample was heated to 800 °C, the formation of the higher boride (AlMgB14) can be clearly seen, but the lower boride (AlMgB4) is still the major phase present. Again, the reaction is not complete, as seen by small peaks for the reactant elements. Moreover, at this temperature, the spinel phase appears. When the sintering temperature is increased to 1000 °C, the higher boride phase is now more abundant than the lower boride, as can be seen from Fig. 6b. At this temperature there is no evidence for the presence of boron and thus the continuation of the increase of AlMgB14 must be at the expense of the lower boride. For this transformation, we propose the following reaction:
7AlMgB4 ¼ 2AlMgB14 þ 5Al þ 5Mg
ð3Þ
This implies that Al and Mg will increase as this transformation takes place. This is supported by the significant increase in the intensity of the Al peaks at 1000 °C, relative to the peaks at 800 °C, as can be seen from Fig. 6b and c. And when the temperature reaches 1250 °C, the higher boride phase, AlMgB14 is the major phase, with the lower boride phase now being a minor phase. However, at this temperature the peaks for Al remain significant. At 1325 °C, the Al peaks decrease, which we attribute to the formation of the spinel phase. This is evident by the increase in the number and the intensity of the spinel peaks. The reaction becomes nearly complete when the temperature of sintering is 1400 °C, as can be seen from Fig. 6c. The evolution of the phase AlMgB14 with temperature, plotted as vol% is shown in Fig. 7. At 1400 °C this phase constitutes 96.2 vol% of the sample. The evolution of composition with sintering temperature is associated with evolution in the microstructure. Fig. 8 shows the changes in the mean grain size for the entire sample with temperature over the range 600–1500 °C. In addition to these results, Table 4 also shows the changes in density and hardness for these samples. The hardness values are reported only for samples with highest densities. The average grain size increased from about 0.4 to 4.8 lm as the sintering temperature increased from 600 to 1500 °C. Over the same range, the density increased from 1.30 to 2.66 g cm3.
Fig. 5. Typical temperature and displacement profiles during the reactive sintering of AlMgB14 at 1500 °C.
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AlMgB14
MgAL2O4
AlMgB4
Al
Mg
β-B
Relative Intensity
a
600°C
800°C
10
20
30
40
50
60
70
2θ (Degrees)
b
AlMgB14
MgA l2O4
AlMgB4
Al
Relative Intensity
Mg
1000°C
1250°C
10
20
30
40
50
60
70
2θ (Degrees)
c
AlMgB14
MgAl2O4
AlMgB4
Al
Relative Intensity
Mg
1325°C
1400°C
10
20
30
40
50
60
70
2θ (Degrees) Fig. 6. X-ray analyses of samples reactively sintered at various temperatures: (a) 600 and 800 °C; (b) 1000 and 1250 °C; and (c) 1325 and 1400 °C.
On the basis of the results shown in Fig. 6, formation of AlNgB14 begins at as low a temperature as 800 °C and is essentially com-
plete at 1325 °C. Thus the prospect of forming nanostructured AlNgB14 is not high since earlier nucleated grains continue to grow
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this work, 3.1 MPa m1/2 is in agreement with previously reported values [23].
4. Conclusions
Fig. 7. Evolution of the phase AlMgB14 as a function of sintering temperature.
The simultaneous synthesis and sintering of AlMgB14 from the elements was investigated by the pulse electric current sintering (PECS) method. The effect of the source and purity of the boron powder was examined. The mechanism of formation of this boride was investigated. The formation of AlNgB14 begins at as low a temperature as 800 °C and is essentially complete at 1325 °C. The results indicate that the prospect of forming nanostructured AlNgB14 is not high since earlier nucleated grains continue to grow, up to the maximum sintering temperature. In addition to phase evolution, the changes in grain size with temperature was also determined. Under optimum conditions, the products contained AlMgB14 with a 93 vol%, the remaining phases contain spinel and impurity from milling. The hardness and fracture toughness of AlMgB14 were determined as 26.1 GPa and 3.1 MPa m1/2, respectively. These values are in agreement with published results. Acknowledgment The financial support of this work by the Army Research Office (ARO) is gratefully acknowledged. References
Fig. 8. Changes in mean grain size with sintering temperature.
up to the maximum sintering temperature. This is seen from the changes of the average grain size with temperature, Table 4, which in view of the increasing fractional content of AlNgB14, reflects primarily the increase in grain size of this boride. Under the conditions utilized in this work, the simultaneous synthesis and consolidation of AlMgB14 can be accomplished at 1400 °C (with a pressure of 123 MPa), giving a product that has the highest content of this boride (96.2 vol%) and the lowest content of the spinel (1.7 vol%). These samples had a hardness of 26.1 GPa and a fracture toughness of 3.1 MPa m1/2. Hardness values of this boride have been reported in literature accounts. In the absence of an intentionally added second phase, the hardness is reported as 28.2 ± 1.8 GPa [16]. As indicated above, hardness values for single crystals of AlMgB14 have been reported to be in the range 20.2–28.6 GPa [12,13]. The value obtained in this work, 26.1 MPa, is consistent with the published results. Higher hardness values were obtained by the addition of a second phase to the boride [16]. And the role of unintentional impurities, those resulting from processing, may explain in part the differences in reported values. From this point of view, the values obtained for single crystals may serve as the reference for the hardness of ‘‘pure” AlMgB14. The fracture toughness of this boride has been shown to be relatively low, having a typical value of 3 MPa m1/2 [16], and recent attempts have been made to improve this value by the addition of a binder alloy [23]. The fracture toughness of the sample made in
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Int. Journal of Refractory Metals & Hard Materials journal homepage: www.elsevier.com/locate/IJRMHM
Mechanism of reactive sintering of MgAlB14 by pulse electric current David J. Roberts 1, Jinfeng Zhao, Zuhair A. Munir * Department of Chemical Engineering and Materials Science, University of California, Davis, CA 95616, USA
a r t i c l e
i n f o
Article history: Received 23 June 2008 Accepted 11 July 2008
Keywords: Aluminum magnesium boride AlMgB14 Ultrahard boride Reactive sintering
a b s t r a c t The reactive sintering of the hard boride AlMgB14 from the elements was investigated using the pulse electric current sintering (PECS) method. The mechanism of formation of AlMgB14 was investigated over the temperature range 600–1400 °C. The formation of AlNgB14 begins at as low a temperature as 800 °C and is essentially complete at 1325 °C. Thus the prospect of forming nanostructured AlNgB14 is not high since earlier nucleated grains continue to grow up to the maximum sintering temperature. Under optimum conditions, the products contained AlMgB14 with a 93 vol%, the remaining phases include spinel and impurity from milling. The product had a hardness of 26.1 GPa and a fracture toughness of 3.1 MPa m1/2. Ó 2008 Elsevier Ltd. All rights reserved.
1. Introduction The search for new hard materials for use in cutting tool, heat sink, abrasive, and coating applications has in the past few decades focused on compounds based on the light elements B, C, and N. Ternary compounds of these elements with a structure similar to diamond are expected to have properties similar to diamond and cubic boron nitride [1,2]. The discovery by Cook et al. [3] of an exceptionally high hardness in AlMgB14 has generated renewed interest in boron-rich compounds as hard materials. These authors reported a hardness of as high as 35 GPa for this boride; with the addition of a second phase, the hardness increases further and approaches the value for cubic BN. Such high hardness, however, is not predicted from theoretical calculations. Lowther [4], using ab initio calculations, showed that the bulk modulus of AlMgB14 is lower than that of B4C, a compound also based on the boron icosahedra structure [5,6], but with a considerably lower hardness (<20 GPa [7]). Using the concept of a ‘‘chemical hardness” [8], the author attributed the hardness difference between AlMgB14 and B4C to the difference in the location of the Fermi level relative to the energy gap [4]. Peters et al. attributed the high hardness of AlMgB14 composites to the intrinsic matrix property and the microstructure [9]. Related to the role of microstructure, observed variation in hardness values were attributed to differences in the powder processing of this boride [10]. The effect of sample preparation on the property of the boride has been recently demonstrated [11]. A significant difference in the bulk modulus
between samples prepared by two methods was attributed to possible differences in impurity incorporation, an observation that underscores the importance of the processing step. The boride AlMgB14 has been prepared by several methods. Okada et al. [12] and Higashi et al. [13] prepared single crystals of AlMgB14 from metal salts and characterized the crystalline structure of the boride and also measured its Vickers hardness, reporting values for the latter in the range 20.2–28.6 GPa. The boride has also been prepared by high-energy milling of the elements followed by hot-pressing or sintering [11]. Takeda et al. [14] used metal hexaborides as precursors for the preparation of AlMgB14 employing the pulsed electric current sintering method (PECS). In addition, thin films of this boride have been synthesized by pulse laser deposition (PLD) [15]. Although the emphasis of the majority of recent investigations on AlMgB14 is on its mechanical properties [10,15,16–18], other investigations have focused on the electrical and thermoelectrical (TE) properties of this boride. Takeda et al. [14] concluded that the boride has little prospect as a TE material because of low electrical conductivity, despite a relatively high Seebeck coefficient. In this paper we present results from an investigation on the simultaneous formation and consolidation of AlMgB14 from the element by the pulse electric current sintering (PECS) method. Takeda et al. [14] used this method, but their starting materials are different and their aim was the investigation of the thermoelectric properties of AlMgB14. 2. Experimental materials and methods
* Corresponding author. E-mail address: [email protected] (Z.A. Munir). 1 Present address: Lawrence Livermore National Laboratory, 7000 East Avenue, Livermore, CA 94551, USA. 0263-4368/$ - see front matter Ó 2008 Elsevier Ltd. All rights reserved. doi:10.1016/j.ijrmhm.2008.07.009
The starting materials were elemental aluminum, magnesium, and boron. The aluminum powder was 99.97% pure (metals basis), 325 mesh (maximum particle size approximately 44 lm)
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obtained from Alfa Aesar (Ward Hill, MA). The magnesium powder was 99.8% pure, 325 mesh, and also obtained from Alfa Aesar. Three types of starting boron were used: amorphous boron powder (99% pure with a sieve classification of 325 mesh), crystalline boron powder (90% pure with an average particle size of <5 lm), and course powder from pulverized boron pieces (<20 mm). All boron materials were obtained from Alfa Aesar. The latter type was used on the premise that the large pieces have smaller amounts of surface oxide contamination. Because oxygen contamination of the preliminary powders was the main concern in producing a single-phase product (i.e., void of the spinel phase MgAl2O4), the storing and handling of powders were done in a glove box under an argon atmosphere. The oxygen level in the glove box was 60.20 ppm. The only case in which the powders were not handled under a controlled atmosphere was during the X-ray diffraction analysis of the starting Al, Mg and B powders and the milled, pre-sintered powders. These powders were initially scanned while sealed in plastic wrap. However, the plastic wrap caused too much interference to give clear XRD patterns of the milled powders. So the plastic was removed and the exposed powders were scanned again in the X-ray diffractometer. From preliminary analysis, it was determined that the source of the oxygen contamination was most likely to be the boron powder. Hence, several different boron-processing methods were tested in an attempt to find the method that yielded a final boride product with the minimum amount of spinel contamination. The aluminum and magnesium powders were used directly from the manufactures containers. The processing methods hereafter will be referred to by the boron type, as can be seen in Table 1. Stoichiometric amounts of various boron powders, aluminum, and magnesium powders (total mass, 7 g) were subjected to high-energy milling in a planetary mill (Fritsch Pulversette 5). The jars were Sialon-lined and the balls were zirconia (10 mm diameter). A charge ratio (ball to powder mass ratio) of 14 was used. The powders were introduced into the jars inside the glove box. The jars were then sealed with parafilm and electrical tape and removed from the glove box. The sealed jars were placed in the planetary mill, which was then operated at a rotation speed of 250 rpm with a milling cycle of 1 h milling, 1 h rest for a total of 20 cycles (40 h total time; 20 h milling time). After the milling was completed, the powders were removed from the jars inside the glove box. In nearly all cases, milling was done in dry condition. However, in a few cases the powder mixture was milled in methanol. For these, the milling time was 8 h. In cases of wet milling, the powders were subsequently dried under vacuum for 72 h. As indicated in Table 1, boron powders (B-3) were annealed in vacuum to reduce the oxygen contamination. In the presence of moisture in the air, boron oxide transforms to the hydroxide
B2 O3 ðsÞ þ 3H2 OðgÞ ¼ 2BðOHÞ3 ðsÞ
having a DG < 0 at 300 °C. The amorphous boron powder (99% pure) was annealed in vacuum inside an alumina crucible that had been sprayed with boron nitride. Annealing was done in a Brew vacuum furnace for 2 h at 1500 °C under an ambient pressure of approximately 6.7 103 Pa. At this temperature, the vapor pressure of B2O3 is 1.88 103 atm. The annealing treatment also resulted in the crystallization of the amorphous boron. Using a similar treatment on amorphous boron with a purity of 99.7 wt%, Brandstötter and Lengauer [19] reported a decrease on oxygen content from 10,000 to 10 ppm after vacuum treatment at the same conditions, 1500 °C for 2 h. Also, as shown in Table 1, the amorphous boron (B-1) was milled in methanol. This was attempted on the basis of an unpublished research indicating that B2O3 reacts with methanol to produce boron esters which volatilize during the subsequent drying of milled powder [20]. The microstructure of the milled powder mixture is shown in Fig. 1. The figure shows agglomerated particles some as large as about 4 lm. X-ray analysis of the milled samples, Fig. 2, showed the presence of the reactant elements only with no indication of product formation during the milling. The figure shows peaks belonging to boron, indicating that milling had resulted in the crystallization of the initial amorphous phase. Grain size analysis on the milled powders was made by the Williamson-Hall method [21]. The result gave grain sizes for Al and Mg of 420 and 46 nm, respectively. In view of the limitation of this method with respect to grain size (100 nm), the value obtained for Al is taken as only approximate. The limited number of boron peaks and their low intensity made the grain size evaluation of this element by this method uncertain. Approximately 2.5 g of the milled powder mixtures were placed in a graphite die which was then placed in a zip-sealed plastic bag and transferred from the glove box to the spark plasma sintering apparatus. The die and plungers used for sintering the milled powders were made of high-density graphite. The die was 3.81 cm high with an outer diameter of 4.45 cm and an inner diameter of 1.91 cm. The plungers were 2.54 cm high and 1.91 cm in diameter. The inside of the die was bored an additional (approximately) 250 lm to accommodate a liner. Graphite foil approximately 125 lm thick was used as a liner inside the die. The liner acted as a barrier to inhibit the sintering powders from bonding to the graphite die. To keep the powders from bonding to the plungers, graphite foil was cut into circles about 1.91 cm in diameter and a double layer was placed between the powders and the plungers.
ð1Þ
Reaction (1) has a DG = 60.9 kJ at 25 °C. And the formed hydroxide is relatively volatile, with the reaction
BðOHÞ3 ðsÞ ¼ BðOHÞ3 ðgÞ
ð2Þ
Table 1 Boron samples and their treatment Boron designation
Source
Treatment
B-1 B-2 B-3
99.0%; amorphous 90%; crystalline 99.0%; amorphous
B-4 B-5
99.0%; amorphous 99.5%; crystalline; crushed pieces
None None Vacuum annealed (1500 °C) Milled in methanol; 8 h None
557
Fig. 1. SEM image of milled reactants Al, Mg, and B.
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The simultaneous synthesis and consolidation of AlMgB14 from powder mixtures was performed in a spark plasma sintering apparatus (SPS). The apparatus is the SPS-1050 made by the Sumitomo Mining Company, Ltd. (Kawasaki, Japan). The power supply has a 15 V, 5000 A capacity. The current generated is a pulsed DC current with a cycle of 12 pulses on and 2 pulses off, with each pulse having a duration of about 3 ms. The uniaxial press can provide up to 100 kN of force. The die containing the powder mixture was taken out of the zip bag and placed in the SPS, which was quickly closed and evacuated. Before sintering, the powders in the die were cold-pressed for approximately 20 s under a load of 10 kN (35 MPa). The load was then reduced to 5 kN (17.5 MPa). Once the SPS chamber pressure reached 10 Pa, sintering was initiated. A current was then provided to increase the temperature in the following manner: from room temperature to 600 °C in 2 min at a pressure of 17.5 MPa, held at 600 °C for 3 min as the pressure was increased to 63 MPa, then from 600 to 1500 °C the temperature was increased at a rate of 100 °C min1. The samples were than held at the final temperature for 5 min, after which the power was turned off and the samples were allowed to cool down. Subsequent experiments were carried out under a higher pressure of 123 MPa. A pyrometer (CHINO, model IR-AH, Japan), focused on the outer wall of the die, was used to measure the temperature of the system. The density of the samples was calculated from geometric and mass determinations and also by the Archimedes method. The microstructure of the samples was determined by optical and scanning electron microscopy. An FEI XL30-SFEG scanning electron microscope (SEM) with energy dispersive X-ray spectroscopy (EDX) capabilities was used to observe the polished surface of the mounted samples. A portion of the unmounted sample was
Fig. 3. SEM image of sample made with annealed boron and sintered at 1500 °C.
broken off and the fractured surface of the sintered samples was also observed under the SEM. Chemical analysis of the surface was done using the EDX function of the SEM. For X-ray analysis a portion of the sample was pulverized. Analysis was made with a Scintag X-ray Diffractometer (XRD), model XDS 2000, with Cu Ka radiation, k = 1.540562 nm. The scanning step size was 0.02 degrees for the quick scan with a hold time of 1.5 s per step, and 0.005 degrees with a hold time of 4.5 s per step for the slow scan. Data from the XRD long scan times were used to
Al Mg
Relative Intensity
β-B B
10
20
30
40
50
60
70
2θ (degrees) Fig. 2. X-ray analysis of milled reactants.
Table 2 Phase analyses (vol%) of samples sintered at 1500 °C with different boron types Sample B-1, B-2, B-3, B-5, B-4, B-1,
Al, Al, Al, Al, Al, Al,
Mg Mg Mg Mg Mg Mg (wet-milled)
AlMgB14
Spinel
Al
Mg
B
AlMgB4
Mg3(BO3)2
ZrB2
Si3N4
85.6 22.4 92.6 29.5 85.1 9.2
6.8 56.8 3.3 22.0 6.3 78.6
1.3 1.6 1.2 3.3 0.7 5.0
0.1 0 0.2 1.0 0 0
3.0 11.6 1.1 8.4 3.5 4.8
0.2 0 0.8 1.0 0 0.4
1.6 4.0 0.4 12.2 2.4 0.2
0.3 0.3 0.2 6.1 1.4 0.4
1.1 3.4 0.4 16.5 0.7 1.5
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Fig. 4. EDX analyses of regions in Fig. 3: (a) The main, dark, phase; (b) gray region; and (c) bright, nearly white region.
Table 4 Grain size, density, and hardness of AlMgB14 samples sintered at various temperatures
Table 3 EDX analyses (at.%) of different regions in Fig. 3
Dark region (Fig. 4a) Off white (Fig. 4b) Bright white (Fig. 4c)
B
O
Mg
Al
Zr
Temperature (°C)
Mean grain Size (lm)
Density (g cm3)
Hardness (GPa)
92.7 0 92.79
0 39.15 0
4.22 22.53 3.29
3.04 38.31 2.50
0 0 1.42
1500 1400 1325 1250 1000 800 600
4.76 3.87 3.44 1.75 1.58 0.37 0.36
2.66 2.68 2.75 2.32 – 1.56 1.39
23.9 26.1 24.6
perform a grain size analysis on the pulverized samples using the Williamson-Hall method [21].
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The Vickers hardness (Hv) was determined, using a 1 kg load and a 15 s dwell time on a Buehler Micromet 2004 microhardness tester. Indents were made in 10 random locations on the sample. The sample fracture toughness was calculated from the length of the cracks formed at the corners of the indents [22]. 3. Results and discussion The initial aim was to determine which of the types of boron provides the best result. Regardless of the type of boron used or the method of milling, all samples sintered at 1500 °C (at 63 MPa) contained the phase AlMgB14, as can be seen from Table 2 which lists representative results for the different conditions. The phase compositions in the table are in vol%, calculated from X-ray data using the PowderCell software. The results provide a wide range of AlMgB14 content, from about 10 vol% to a high of about 93 vol%. They indicate that wet milling of all reactants (in methanol) results in samples with the least amount of the desired boride, probably because of ineffective mixing and folding of the reactant metals and the possible oxygen contamination from water in the alcohol. The sample that gave the highest boride content and least amount of spinel was that in which annealed boron (B-3) was used. SEM image of a sample made with annealed boron and reactively sintered at 1500 °C under a pressure of 63 MPa is shown in Fig. 3. It should be recalled this sample contained the highest vol% of the desired boride, 93%. EDX analysis on this sample showed the main phase (dark region) contains Al, Mg, and B, Fig. 4a, and thus represents the desired boride phase. The inclusions in Fig. 3 are represented by two regions: a relatively major one with an off-white appearance and a minor region with bright white appearance. EDX analysis of the former, Fig. 4b, indicates that this is the spinel phase, while the latter, Fig. 4c, is a zirconium-rich phase, possibly a boride originating from contamination from the milling balls. Because of the small size of these latter contaminants, the EDX analysis includes peaks of the background, the boride phase. The results of the EDX analysis are presented in Table 3. As stated above, the vol% of AlMgB14 in this sample was about 93%, which would mean that the other phases (including the spinel) constitute about 20% on the basis of area, as depicted in the SEM image (Fig. 3). The average density of the samples providing the best results from Table 4 was determined as 2.66 and 2.62 by the geometric and Archimedes methods, respectively. The theoretical density of AlMgB14 is reported as 2.6 g cm3. To understand the process of formation and densification of AlMgB14 from the elements we conducted SPS experiments to intermediate temperatures. Typical SPS temperature and displacement profiles for a sample using annealed boron are shown in Fig. 5. The displacement curve reflects uniaxial changes in the sample dimension and thus relate to shrinkage; the displacement is positive when the sample shrinks and negative when it expands. The lowest temperature measured by the pyrometer is 600 °C and thus the presence of the flat portion on the initial temperature profile, ending at point 1a in Fig. 5. The increase in shrinkage at point 1b is the result of increasing the applied pressure from 17.5 to 63 MPa. On the other hand, the considerable shrinkage beginning at point 2b (corresponding to 1250 °C) is due to the simultaneous synthesis and densification of the boride. The temperatures indicated along the displacement curve are those used to halt the process for investigation of the synthesis and consolidation process. Samples were heated along the temperature curve of Fig. 5 and then the heating is stopped at selected temperatures of 600, 800, 1000, 1250, 1325 and 1400 °C. For the case of 1325 and 1400 °C, the pressure utilized was 123 MPa. X-ray analyses were then performed on the samples to determine phase composition. Fig. 6a
shows the results for 600 and 800 °C, Fig. 6b shows the results for 1000 and 1250 °C, and Fig. 6c shows the results for 1325 and 1400 °C. As can be seen from Fig. 6a, when the milled powders are heated in the SPS to 600 °C, partial reaction of the elements takes place, producing the lower boride phase, AlMgB4. No evidence of the AlMgB14 can be seen at this temperature. At this temperature, the reaction to form any boride is not complete, as seen by the presence of peaks for Al and Mg. However, when the sample was heated to 800 °C, the formation of the higher boride (AlMgB14) can be clearly seen, but the lower boride (AlMgB4) is still the major phase present. Again, the reaction is not complete, as seen by small peaks for the reactant elements. Moreover, at this temperature, the spinel phase appears. When the sintering temperature is increased to 1000 °C, the higher boride phase is now more abundant than the lower boride, as can be seen from Fig. 6b. At this temperature there is no evidence for the presence of boron and thus the continuation of the increase of AlMgB14 must be at the expense of the lower boride. For this transformation, we propose the following reaction:
7AlMgB4 ¼ 2AlMgB14 þ 5Al þ 5Mg
ð3Þ
This implies that Al and Mg will increase as this transformation takes place. This is supported by the significant increase in the intensity of the Al peaks at 1000 °C, relative to the peaks at 800 °C, as can be seen from Fig. 6b and c. And when the temperature reaches 1250 °C, the higher boride phase, AlMgB14 is the major phase, with the lower boride phase now being a minor phase. However, at this temperature the peaks for Al remain significant. At 1325 °C, the Al peaks decrease, which we attribute to the formation of the spinel phase. This is evident by the increase in the number and the intensity of the spinel peaks. The reaction becomes nearly complete when the temperature of sintering is 1400 °C, as can be seen from Fig. 6c. The evolution of the phase AlMgB14 with temperature, plotted as vol% is shown in Fig. 7. At 1400 °C this phase constitutes 96.2 vol% of the sample. The evolution of composition with sintering temperature is associated with evolution in the microstructure. Fig. 8 shows the changes in the mean grain size for the entire sample with temperature over the range 600–1500 °C. In addition to these results, Table 4 also shows the changes in density and hardness for these samples. The hardness values are reported only for samples with highest densities. The average grain size increased from about 0.4 to 4.8 lm as the sintering temperature increased from 600 to 1500 °C. Over the same range, the density increased from 1.30 to 2.66 g cm3.
Fig. 5. Typical temperature and displacement profiles during the reactive sintering of AlMgB14 at 1500 °C.
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AlMgB14
MgAL2O4
AlMgB4
Al
Mg
β-B
Relative Intensity
a
600°C
800°C
10
20
30
40
50
60
70
2θ (Degrees)
b
AlMgB14
MgA l2O4
AlMgB4
Al
Relative Intensity
Mg
1000°C
1250°C
10
20
30
40
50
60
70
2θ (Degrees)
c
AlMgB14
MgAl2O4
AlMgB4
Al
Relative Intensity
Mg
1325°C
1400°C
10
20
30
40
50
60
70
2θ (Degrees) Fig. 6. X-ray analyses of samples reactively sintered at various temperatures: (a) 600 and 800 °C; (b) 1000 and 1250 °C; and (c) 1325 and 1400 °C.
On the basis of the results shown in Fig. 6, formation of AlNgB14 begins at as low a temperature as 800 °C and is essentially com-
plete at 1325 °C. Thus the prospect of forming nanostructured AlNgB14 is not high since earlier nucleated grains continue to grow
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this work, 3.1 MPa m1/2 is in agreement with previously reported values [23].
4. Conclusions
Fig. 7. Evolution of the phase AlMgB14 as a function of sintering temperature.
The simultaneous synthesis and sintering of AlMgB14 from the elements was investigated by the pulse electric current sintering (PECS) method. The effect of the source and purity of the boron powder was examined. The mechanism of formation of this boride was investigated. The formation of AlNgB14 begins at as low a temperature as 800 °C and is essentially complete at 1325 °C. The results indicate that the prospect of forming nanostructured AlNgB14 is not high since earlier nucleated grains continue to grow, up to the maximum sintering temperature. In addition to phase evolution, the changes in grain size with temperature was also determined. Under optimum conditions, the products contained AlMgB14 with a 93 vol%, the remaining phases contain spinel and impurity from milling. The hardness and fracture toughness of AlMgB14 were determined as 26.1 GPa and 3.1 MPa m1/2, respectively. These values are in agreement with published results. Acknowledgment The financial support of this work by the Army Research Office (ARO) is gratefully acknowledged. References
Fig. 8. Changes in mean grain size with sintering temperature.
up to the maximum sintering temperature. This is seen from the changes of the average grain size with temperature, Table 4, which in view of the increasing fractional content of AlNgB14, reflects primarily the increase in grain size of this boride. Under the conditions utilized in this work, the simultaneous synthesis and consolidation of AlMgB14 can be accomplished at 1400 °C (with a pressure of 123 MPa), giving a product that has the highest content of this boride (96.2 vol%) and the lowest content of the spinel (1.7 vol%). These samples had a hardness of 26.1 GPa and a fracture toughness of 3.1 MPa m1/2. Hardness values of this boride have been reported in literature accounts. In the absence of an intentionally added second phase, the hardness is reported as 28.2 ± 1.8 GPa [16]. As indicated above, hardness values for single crystals of AlMgB14 have been reported to be in the range 20.2–28.6 GPa [12,13]. The value obtained in this work, 26.1 MPa, is consistent with the published results. Higher hardness values were obtained by the addition of a second phase to the boride [16]. And the role of unintentional impurities, those resulting from processing, may explain in part the differences in reported values. From this point of view, the values obtained for single crystals may serve as the reference for the hardness of ‘‘pure” AlMgB14. The fracture toughness of this boride has been shown to be relatively low, having a typical value of 3 MPa m1/2 [16], and recent attempts have been made to improve this value by the addition of a binder alloy [23]. The fracture toughness of the sample made in
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