Processing and mechanical properties of pure mg and in situ aln reinforced mg 5al composite 3

During the ball milling process, the mechanical energies transferred to the powder mixture in the presence of highly reactive transition metal surfaces accelerates the decomposition of p

Commercial AlN powder is usually produced by carbothermal reduction of Al2O3 or Al(OH)3 and direct nitridation of aluminium [1,9-11] In the carbothermal method, the starting material is reduced by carbon and reacted with N2 at high temperatures (1300°C) In the direct nitridation method, Al metal powder reacts with N2 or NH3 at high temperatures (1150°C) It is often prepared by methods such as DC arc-discharge [12], sublimation recrystallization [13,14], sodium flux [15], volatilization-condensation of AlN powder [16], chemical vapor deposition [17] and combustion synthesis or self-propagating high-temperature synthesis [18, 19]

MM technique has been mainly used for preparation of materials which are difficult to

be prepared by the ordinary melting methods, for example, oxide dispersion strengthened alloys and amorphous or supersaturated solid solution alloys with

remarkably extended solid solubilities This method has been extended to use for low

temperature synthesis of AlN by direct nitridation of Al metal in nitrogen and dry

ammonia atmosphere under pressure of 600 kPa through mechanochemical reaction

[20,21] Since the formation of AlN from Al metal occurs by an exothermic reaction, it

proceeds spontaneously at higher temperatures due to the heat generated by the

nitridation [22]

(a) (b) Figure 3.1 (a) Morphology of Al powder and (b) structure of pyrazine

In this study, in-situ formed AlN composite powder is obtained by milling the mixture

of elemental Al with pyrazine (C4H4N2) which shows a ring-type organic material as

shown in Figs 3.1(a) and (b) Due to the high concentration of nitrogen in parazine,

relatively fast direct nitriding reaction compared to NH3 gas can be maintained during

the milling process During the ball milling process, the mechanical energies

transferred to the powder mixture in the presence of highly reactive transition metal

surfaces accelerates the decomposition of pyrazine into simple HxC-, HxN- or CN-

radicals and side chains [23] The dissociated nitrogen (N) from pyrazine molecules is

expected to react with highly reactive Al metal surfaces especially when Al is in nanocrystalline structure after certain period of MM

3.2 Experimental

Al powder (Alfa Aesar, -325 mesh, 99.5% purity) and pyrazine (H4C4N2) (Alfa Aesar, 99+% purity) were used as the starting materials in this study 14g Al powder mixed with pyrazine in the amount to satisfy the stoichiometric ratio of Al:N=1:1 was loaded into 250 ml stainless steel vial together with 20 carbon steel balls in a 99.9% pure argon atmosphere in an AMBRUAN glove box A Retsch PM100 Planetary Ball Mill was employed for MM at 300 rpm The weight ratio of Al powder to ball was 1:20 For blended powder mixture, Al and pyrazine were thoroughly mixed in an agate mortar with an agate pestle

After every 20h of milling, a small quantity of powder was withdrawn and annealed at

500 and 1250°C in a tubular furnace for one hour in purified argon flow The blended powder mixture was also annealed for the comparison study An X-ray diffraction (XRD) measurement using Cu Kα radiation operating at 40 kV and 30 mA was carried out for structural examination of MMed powders and MMed powders after annealing

at 500°C and 1250°C (designated hereafter as MMed, MMed-500 and MMed-1250 samples respectively, where xx is milling hours and the blended powder mixture is indicated as 0h) The microstructures of the samples were examined using a Quanta 200F field emission scanning electron microscope (FESEM) and Jeol 2010F TEM Shimadzu DTG-60/60H was employed for the simultaneous measurements of thermogravimetry and differential thermal analysis

xxh-3.3 Results and discussion

Figure 3.2 X-ray diffraction patterns of 0h-MMed Al-Pyrazine mixture after annealing

at 500°C and 1250°C for 1h

Fig 3.2 shows the X-ray diffraction patterns of 0h-MMed Al-pyrazine mixture annealed at 500 and 1250°C for one hour in purified argon gas flow Only Al peaks are detected in the 0h-MMed-500 sample indicating no reaction between Al and pyrazine For the 0h-MMed-1250 sample, significant oxide formation is detected from XRD patterns with high intensity alumina (Al2O3) peaks Although annealing was carried out in high purity argon gas, oxidation could not be fully prevented Since the melting temperature of pyrazine is 54°C which is far below the melting temperature of Al, it evaporated at much lower temperature than that of reaction between Al and N Hence, AlN could not be formed

The XRD patterns of the MMed aluminium-pyrazine mixtures at different milling durations are shown in Fig 3.3 All the diffraction peaks from the 20h- and 30h-MMed samples are from the pure Al implying no formation of compound between N/carbon

(C) atoms and Al The DTA traces of MMed powder for different milling durations in Fig 3.4 show endothermic peaks at 668°C due to Al melting from the powders milled for 20h and 30h It is possible that the mechanical energy supplied to the milling process was not high enough to break the bonds between C and N in the pyrazine molecules and/or to overcome the diffusion barrier for N atoms to diffuse into the Al particles

40h 60h 80h

Figure 3.3 X-ray diffraction patterns of MMed Al-Pyrazine powder at different milling durations

As milling progressed, N atoms gradually diffused into highly defective Al particles After 40h of milling, AlN started to form as clearly shown in Fig 3.3 Although weak

Al peaks are detected, the amount of residual unreacted Al may not be large enough to produce any endothermic peak corresponding to the Al melting in DTA traces (Fig 3.4) It can be seen that the AlN peaks are broad, which corresponds to the characteristics of nanocrystallites with the presence of interfacial components composed of Al, H and N atoms [20] The average crystalline sizes estimated using Scherrer’s formula based on the theory of broadening of XRD diffraction peaks are 10,

9, 8 and 8 nm for 40, 60, 80 and 100h-MMed powders respectively

0 0

30h20h

Figure 3.4 DTA traces of the Al-Pyrazine powders MMed for different milling durations

Due to low melting point and asymmetric structure of pyrazine, it is unstable under thermal treatment and/or mechanical activation [23] Pyrazine therefore decomposed to some extent after 40h of milling and a certain amount of released N reacted with Al with the evidence of weak and broad AlN peaks The proportion of AlN in the powder increased with further milling Higher and distinct AlN peaks indicating higher

crystallinity of AlN particles could be observed in the 100h-MMed sample Disappearance of endothermic peaks at 668°C in DTA traces corresponding to Al melting in the 60, 80 and 100h-MMed samples confirms the complete transformation

of Al to AlN and Al2O3

Al2O3 oxide peaks were detected in the MMed powders except in 20h- and 30h-MMed powder samples The sources of Al2O3 formation in the AlN powder could have originated from the oxide layers on the Al raw material [24], possible oxidation during milling and powder handling due to its highly pyrophoric nature

In the dry-milling process, under continuous impact and shear stress during MM,

Al2O3 passivation layer on the as-received Al particles was fractured creating new surfaces to react with free N dissociated from pyrazine ring structure Due to the strong affinity between Al and N, the N atoms were adsorbed on the newly created Al surfaces and then incorporated in the interfaces by the pressure welding of the Al powder followed by diffusing into the Al matrices through grain boundaries, dislocations and other defects [25] Milling induced dislocation and vacancy density in the MMed Al enhanced the favorable situation for direct chemical nitridation following the decomposition of pyrazine molecules During mechanochemical activation, increased diffusion of nitrogen was attributed to development of nanostructural or amorphous phase [26] The exothermic behavior owing to the extensive heat of reaction enhances the reaction rate [27]

Figs 3.5(a) and (b) show the XRD patterns of MMed-500 and MMed-1250 samples respectively The XRD patterns of MMed-500 samples are very much similar to those

of as-milled samples indicating the negligible effect of annealing at 500°C No AlN structure could be detected in 20h and 30h-MMed-500 samples After annealing at 1250°C, AlN was found to form in both 20h and 30h-MMed powders In addition to AlN formation, oxide and carbide were detected from Al2O3 and Al4C3 diffraction peaks Al4C3 is the only intermediate compound in Al-C binary phase diagram [28] A weak peak which might correspond to FeAl2 was observed in all MMed-1250 samples

Al and Fe contamination from milling media formed FeAl2 intermetallic during milling Compared to the powder annealed at 500°C, oxidation was more pronounced and Fe contamination from milling media was clearly identified Average crystalline sizes of AlN after annealing at 1250°C were estimated to be 14, 13, 6, 9, 10 and 10nm for 20, 30, 40, 60, 80 and 100h-MMed samples respectively indicating no significant grain growth during annealing

Table 3.1 Contents of C, H, N, Al and Fe under different conditions

Sample C (wt.%) H (wt %) N (wt %) Al (wt %) Fe (wt %)

80h-MMed-1250 28.16 <0.50 16.22 50.31 1.12 60h-MMed-1250 27.62 2.51 13.02 34.95 0.45 40h-MMed-1250 28.62 0.77 15.78 40.96 0.58 30h-MMed-1250 22.34 <0.50 11.26 45.35 0.86 20h-MMed-1250 17.58 0.51 8.92 52.87 0.84 100h-MMed-500 27.59 1.36 16.09 42.20 1.07 80h-MMed-500 26.88 1.36 14.18 41.08 0.79 60h-MMed-500 27.03 1.35 14.32 45.21 0.90 40h-MMed-500 26.41 2.33 13.49 43.73 0.73 30h-MMed-500 17.92 1.56 7.80 41.08 0.53 20h-MMed-500 15.97 1.37 7.34 50.55 0.65

From the results of chemical analysis as tabulated in Table 3.1, a certain amount of carbon did exist in the MMed powders However, only the 20h- and 30h-MMed-1250 powders showed the Al4C3 carbide peak but not in 500°C annealed samples After 20h and 30h MM followed by subsequent annealing at 500°C for one hour, the mechanical activation and thermal activation might not be sufficient for carbide formation For longer milling duration, the disappearance of carbide peaks might be due to the emergence of carbide peaks with broadened AlN peaks due to grain refinement with

longer milling duration and the carbide particles are too fine to produce diffraction peaks

3.3.2 Formation of alumina and AlN whiskers

Various microstructures of AlN have been reported in the literature such as fibers, agglomerated particles, coarse granules, hexagonal whiskers, and faceted particles In the synthesized AlN products, different morphologies coexisted in most cases [29] In this study, the Al powders were subjected to micro-forging, fracture, agglomeration, and deagglomeration during MM process and the crystals in the deformed powder particles were heavily stressed and strained in a rather inhomogeneous manner [30] It

is generally accepted that formation of protrusions or whiskers results from local stress-relaxation process [31-33] Since the stress and strain distribution in the MMed

Al powder is inhomogeneous and the fluctuations of growth conditions such as localized distribution of absorbed nitrogen molecules, degree of Al vapor saturation and reaction temperature, different morphologies of AlN whiskers and particles coexisted in this present material

After annealing Al-Pyrazine blend at 1250°C for one hour, white cotton-like Al2O3

whiskers were observed to be surrounding the spherical shape Al melt in the middle TEM image with selected area electron diffraction (SAED) pattern (insert) in Fig 3.6(a) reveals that the whiskers are alumina single crystal confirmed by the lattice spacing of 3.5Å corresponding to diffraction from the (012) plane Fig 3.6(b), a high resolution TEM (HRTEM) image of the lattice pattern from 0h-MMed-1250 sample, shows an inter-planar spacing of 2.559Å corresponding to the (104) planes of the

Al2O3 crystal which appears to grow in a direction angled at ~30° with normal to the (104) plane

Figure 3.6 TEM images of (a) alumina whisker and SAED pattern (insert) and (b) HRTEM image of lattice pattern from the same whisker (growth direction along black arrow)

Fig 3.7 shows the morphologies of 0h-MMed-1250 sample In entangled alumina whiskers, network-like structures with a radial outgrowth from a single stem-like structure can be seen in Fig 3.7(a) The alumina whisker has a (001) facet (C-facet) that appeared as a flat section on the top with a simple six-sided structure as shown in Figs 3.7(b), (c), (d) and (e) Lack of droplet on the tip of the whisker excluded a VLS mechanism and the possible growth mechanism of the alumina could be vapor-solid (VS) route From Figs 3.7(d) and (e), a repetitive stacking of the platelets along the growth direction gives rise to the complete structure However, the existence of such distinct layers disappeared while smooth appearance was significant with pillar-like structure in Fig 3.7(c) Such structure was resulted due to adatom hopping as a result

of a reducing Ehrlich-Schwoebel (ES) barrier (step-edge barrier which accounts for the diffusion barrier of an adatom over the edge of a step) and lateral growth became more pronounced to generate near perfect uniform dimension [34] When adatoms can easily overcome the ES barrier and get down to the step-edge, layer-by-layer growth is promoted Multilayer growth mode starts to appear if there is an asymmetry in the incorporation rate of adatom into the step and it is much lower than surface diffusion [35] Fig 3.7(d) shows the incomplete crystal growth of alumina pillar with hexagonal flat top oozing out from the surface

Fig 3.7(f) shows whiskers with kinks resulting from sudden change in their growth orientation It is unlikely due to whisker collapsing by their own weight since straight whiskers up to 500 μm could be observed It is possible that kinking may be caused by the uneven flow of material across the growth interface [33]

Fig 3.8 shows the FESEM images of MMed Al-Pyrazine mixture at different milling durations ranging from 20h to 100h Since Al is ductile material, the particles were heavily agglomerated during MM and the aggregates were composed of particle size ranging from micron to submicron As reported by William [36], a broad particle size distribution and the presence of large aggregates were observed Refinement and reduction in particle size with prolonged milling is evident in Fig 3.8

No whisker formation could be detected in all MMed samples annealed at 500°C This

is because the annealing temperature is well below the melting temperature of Al to have reaction between Al and N The samples annealed at 1250°C consisted of both

100h-AlN particles and whiskers as shown in Fig 3.9 Figs 3.10 (a) and (b) show a TEM image with SAED pattern (insert) and HRTEM image of lattice pattern from 100h- MMed-1250 sample It reveals that the AlN particles are poly-crystalline with an inter-planar spacing of 2.74 Å corresponding to the (100) planes of the AlN crystal grown along the [100] direction