СТРУКТУРА, ФАЗОВЫЕ ПРЕВРАЩЕНИЯ И ДИФФУЗИЯ
УДК 669.71 '24:539.89
EFFECT OF MILLING PARAMETERS ON THE SYNTHESIS OF Al—Ni INTERMETALLIC COMPOUND PREPARED BY MECHANICAL ALLOYING
© 2015 г. Kahtan S. Mohammad*, Haider T. Naeem*, **
*School of Materials Engineering, University Malaysia Perils, 02600 Jejawi, Perlis, Malaysia **College of Engineering, Almuthana University, Almuthana, Iraq e-mail: haider_neem@yahoo.com
Поступила в редакцию 28.07.2014 г.; в окончательном варианте — 15.09.2014 г.
Nanocrystalline Al—Ni intermetallic compounds were synthesized with different percentages of nickel by mechanical alloying (MA) of elemental powders. In all MA runs, the ball-to-powder weight ratio was 10 : 1, the rotation speed was 350 rpm, and the milling time ranged from 4 to 12 h. The phase evolution and microstructural changes of the powders during MA were investigated by X-ray diffraction (XRD), scanning electron microscopy, and energy-dispersive X-ray spectroscopy analyses. The crystallite sizes of milled powders were estimated based on the broadening of XRD peaks using the Williamson-Hall formula. Results showed that an optimum MA time of 12 h led to the formation of solid solutions of Al—Zn—Mg and Ni, which can be added to many Al—Ni intermetallic compounds. Furthermore, an Al-Ni intermetallic phase with <20 nm crystallite size was obtained.
Keywords: intermetallics, mechanical alloying, nanocrystalline materials, SEM.
DOI: 10.7868/S0015323015090077
1. INTRODUCTION
Many intermetallic compounds show an interesting combination of physical and mechanical properties [1]. The technological "pull" comes mainly from the needs of the aerospace industry for new high-temperature structural materials with properties that cannot be met by ceramics or by conventional superalloys. Intermetallic compounds have higher melting temperatures than superalloys and better toughness than ceramics because of metallic bonding [2]. Al—Ni intermetallic compounds have been the focus of considerable interest and are considered potential materials for high-temperature and coating applications because of their low densities, high thermal conductivity, and high corrosion and oxidation resistance at high temperatures [2, 3]. In the case of intermetallic compounds, where Al constitutes majority of the alloy, low density is an additional advantage. Melting and casting is a frequently used method for the fabrication of nickel aluminides [4]. However, this method has some disadvantages, which resulted from the large difference between the melting points ofAl and Ni were connected with evaporation and oxidation [4, 5]. Alternatively, some novel processing techniques, including the Exo-Melt process [5], self-propagating high-temperature synthesis [6], friction stir processing [7, 8], and mechanical al-
loying (MA) [9], have been used to produce nickel aluminide intermetallic compounds. During the MA process, alloys are formed by solid-state reaction. Therefore, the MA technique overcomes problems, such as the large difference in melting points of the alloying components and evaporation or segregation that could occur during melting and casting. MA is a dry powder processing technique and has been used to synthesize nanostructured intermetallic compounds with ductile behavior. This technique involves repeated cold welding and fracturing of powder particles in a high-energy ball mill [10]. Cold welding and fracturing allow powder particles to be in contact with each other with clean surfaces and minimized diffusion distance. Effective parameters in MA are time, speed of milling, ball-to-powder weight ratio, and atmosphere [10, 11]. To obtain the optimum characteristics of MA intermetallic powder, the process parameters should be selected properly. MA time, as well as the interfacial reaction between matrix and reinforcement, is one of the important parameters affecting the properties of composites. Therefore, this study aims to investigate the feasibility of producing the Al—Ni intermetallic compounds by MA with different amounts of nickel, as well as different milling times. The phase, micro-
Specification of materials
Powder Description Particle sizes, ^m Purity, % Source
Al Flake D50 of 51 98.00 (stabilized of 2%) Merck KGaA
Zn Rounded D50 of 18 96.00 Merck KGaA
Mg Rounded D50 of 115 98.00 Merck KGaA
Cu Irregular D50 of 39 99.50 Merck KGaA
Ni Rounded D50 0f 11 99.50 Merck KGaA
structural, structural, and microanalysis evolutions occurring in the MA process were also investigated.
2. EXPERIMENTAL PROCEDURE
To synthesize Al milled powders using the MA process, the elemental powder precursors of Al, Zn, and Ni were used as the raw materials for the preparation of two types of, Cu , Mg, milled powders, namely A (Al—5% Zn—2.5% Mg—1.5% Cu—3% Ni) and B (Al—5% Zn—2.5% Mg—1.5% Cu—5% Ni). All compositions are expressed in weight percentage ( wt %). Table 1 lists the purity levels and the particle sizes of the raw materials.
Ball milling of the powder was conducted in a planetary high-energy ball mill (Fritsch model Pulverisette 5) under argon atmosphere using stainless steel balls of 10 to 20 mm diameter. The balls/powder weight ratio was 10 : 1 and the rotation speed was 350 rpm. MA process for the powder mixture was completed at different milling times of approximately 4, 8, and 12 h. Ball milling experiments were periodically halted every 15 min and then resumed for 45 min to prevent a significant increase in temperature. The phases of milled powders A and B were examined via X-ray diffraction (XRD) analysis using a D2 PHASER diffractometer with Cu Ka radiation (X = 1.54184 Â) at 30 kV and 10 mA.
The XRD patterns were recorded within the 2° range of 20° to 80° with a step size of 0.02 and an effective total time of 651.2 s. The crystallite size and lattice strain of the powders were estimated via the Williamson—Hall method, as follows [12]:
$kki cos QhU = jX + 2^Vs~2sin Qhkl, (1)
where cosQhkl is the Bragg diffraction angle, K is the shape factor (0.9), D is the crystallite size, s is the lattice strain, and X is the X-ray wavelength. The instru-
mental corrected broadening was approximated by a Gaussian fit as the full width at half maximum (FWHM), which as calculated using an XRD analysis soft are (X'Pert HighScore) based on each diffraction angle of 20 Morphology of milled po der particles as observed by scanning electron microscopy (SEM) using JEOL JSM-6460LA. Chemical composition was verified by energy-dispersive X-ray spectroscopy (EDS).
3. RESULTS AND DISCUSSION
The initial morphologies of as-mixed and milled powder A are shown in Fig. 1. The as-mixed powder A is flake-like as aluminum powder conglomerates and the remaining powders separate (Fig. 1a). After 4 h of milling, the particles were severely deformed plastically by MA and underwent shape changes from flakelike to semispherical (Fig. 1b). The Ni particle is embedded into the large particles by cold welding. Fig. 1c shows the EDS spectra.
In the SEM morphology of milled powder A at a milling time of 8 h (Fig. 2a), the milled powder particles agglomerate with more diffused nickel particles in the soft aluminum matrix. Large aggregates were formed because of high ductility [13]. The EDS spectrum of the labeled region in Fig. 2b shows the increasing wt % of nickel. At a milling time of 12 h, milled powder A, as shown in Fig. 2c, has small particles and agglomerates, indicating that milling approached the steady-state condition. We also observed that the nanoparticles of nickel are adherent, with scattered intensity.
The EDS analysis showed the increasing wt % of nickel in the aluminum matrix (Fig. 2d). When MA is applied, the particles undergo work hardening and fracture by a fatigue failure mechanism [13, 14].
The morphology of milled powder B at a processing time of 4 h, as shown in Fig. 3a, indicates a uniform
(a)
(b)
3000 2700 2400 2100 1800 1500 1200 900 600 300 0
001
(c)
Al
Zn _Zn Cu
_ I
Ni _Ni|M
OC C
ZAF Method Standardless Quantitativa Analysis
Fitting Coefficient: 0.5636
Element (keV) mass % Error % At %
C K* 0.277 6.90 4.31 14.73
O K* 0.525 5.68 2.24 9.10
"MgK* 1.253 3.45 1.04 3.63
Al K* 1.486 70.85 1.08 67.26
Ni K* 7.471 2.72 5.84 1.19
Cu K* 8.040 1.55 8.02 0.79
Zn K* 8.630 8.43 10.33 3.30
Total 100.00 100.00
Cu Zn
Ni
Ni Cu Zn
liil
123456789 10 keV
Fig. 1. (a) and (b) show the SEM morphology of as-mixed and milled powder A and (c) the EDS microanalysis of milled powder A at a milling time of 4 h.
distribution of the nanoparticles of Ni embedded in the Al—Zn—Mg—Cu cluster. The EDS spectrum of the labeled region in Fig. 3b shows a relative increase in the wt % of nickel in mixture B of approximately 5%. Fig. 3c shows the SEM micrograph of milled powder B after 8 h of milling and the reduction in particle size (compared with Fig. 3a). With a long milling time, particles were broken into smaller pieces [15]. By con-
trast, fragments generated by MA continued to decrease in size in the absence of strong agglomerating forces [16]. We noted the extension of the dissolution of nickel in the aluminum mixture. Fig. 3d shows the EDS analysis of the elements of milled powder B. The SEM micrograph of milled powder B after 12 h of milling, as shown in Fig. 3e, indicates that, as milling time increases, the fragmentation and cold welding by
(a)
006
(b)
(c)
3000 2700 2400 2100 1800
g 1500
1200 900 600 300
0
3000 2700 2400 2100 1800
M
3 1500
o
C 1200
900
600
300
- Zn
I
Zn
Cu |-Cu Ni
NiMg
CO i
I
ZAF Method Standardless Quantitativa Analysis .Fitting Coefficient: 0.5600 Element (keV) mass % Error %At % Al C K 0.277 14.49 3.04 29.61 O K 0.525 2.93
Mg K 1.253 2.89 Al K 1.486 61.40
NiK 1.472 4.73
Cu K 8.040 2.99
Zn K 8.630 10.52
Total 100.00
Ü
1.95 4.50 0.91 2.92 0.92 55.87 4.52 2.00 6.26 1.16 8.07 3.95 100.00
Cu Zn
Ni Ni Cu Zn <
1234567 keV (d)
9 10
008
- Zn
_ Zn
Cu " Cu Ni
ZAF Method Standardless Quantitativa
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