Amorphization in the AlC system by mechanical alloying

Jouraal of

~ S A~D C O M P O U ~ ELSEVIER

Journal of Alloys and Compounds 260 (1997) 121-126

Amorphization in the A1-C system by mechanical alloying N . Q . W u * , J . M . Wu, G . - X . W a n g , Z.Z. Li Department of Materials Science and Engineering, Zhefiang University, Hangzhou 310027 31(k027, P.R. China

Received 3 January 1997; received in revised form 10 March 1997

Abstract

Mechanical alloying of a powder mixture of elemental A! and graphite has been performed in a high-energy ball mill. The structural evolution has been characterized by X-ray diffraction and transmission electron microscopy. The carbide AI4C.~ is first formed as an intermediate product. Further milling leads to destabilization of A!4C3. It is proposed that destai ~lization of AI4C3 is induced by the accumulated defects and the high pressure due ~o collision of the balls Bailing milling of the elemental AI-C powder mixture finally results in a f.c.c, solid solution with a carbon content up to 23 at%. Whereas an amorphous phase is formed in the composition range of 28-50 at% C. © 1997 Elsevier Science S.A. Keywords: Mechanical alloying; AI-C system; Amorphization: Solid state reaction; Phase transformation

1. Introduction Mechanicat alloying (MAI of a bh~ary mixture of elemental powders has been widely used for producing amorphous alloys in various systems [1,2], such as metalmetal systems and transition metal-metalloid systems (e.g. Fe-B [3], Co-B [4]). However, it is known tha~ amorphous afloys of metal-carbon systems are difficult to manufacture by mechanical alloying of the elemental powder m~xtures [51. Ball milling of a M - C (M=Ni, Co) mixture blend has resulted in the supersaturated solid state solution or metastable phase M3C inst3ad of in a amorphous phase [6]. During MA of elemental Fe and. C powders, the partial formed amorphous phase has bee~ detected at the intermediate milling stage. On further ball milling the materials are fully transformed into the metastable carbides [7]. In the present work, the amorphous phase has been obtained by MA of a powder mi×ture of elemental aluminum and graphite. The choice of the Al-C system rests on the fact that the AI-AI4C 3 composite fabricated by MA has a special advantage of good mechanical properties retained at elevated temperatures in addition to excellent room temperature properties [8,9]. However, phase transformation associated with thermodynamics during MA of AI-C powder mixture have not been well understood. In addition, different mechanisms of amorphization in *Corresponding author. 0925-83881971517.00 © 1997 Elsevier Science S.A. All rights reserved. PII S0925-8388(97)00138-2

the different systems during mechanical ~loying have been proposed, such as local melting of materials followed by rapid solidification [10], solid-state reaction similar to the thin fi',m diffusion couples [i1,12], or polymorphous melting transition [13,14]. However, in the Ni-Zr system, Weeber et al. [1,15] have found that the c~3,stalline intermetallic intermediate product is firstly formed and further milling resulted in the transformation of the intermediate product into a homogeneous amorphous alloy. Therefore, the purpose of this paper is to investigate the microstructural development and phase formation during ball milling of the AI-C powder blends. The mechanism of amorphization wdl be discussed.

2. E,:per~mental procedure High purity (99.5 wt%) elemental aluminum (particle size 100 ~in) and graphite powders (particle size 50 p.m) were mixed together to give the desired compositions of Al~oo_~C~ (x= 16, 23, 28, 42.9, and 50 at%). The powder mixtures were then put into steel vials in a glove box under argon atmosphere. Ball milling was pe;formed in a planetary type ball mill with a ball-to-powder weight ratio of 30:1. The milling process was interrupted after periodicA milling periods to take out a small amount of as-milled powders. The as-milled powders were characterized using D~,~V,, V .... a~¢~C~,-,,,,-~,,r (XRD) with Cu Ko~ r~mation and a Philips EM420 transmission electron micro-

N.Q. I~, el al. I Journal of Alloys and Compounds 260 (1997) 121-~26

122

scope (TEM). The powders for TEM examinatmn were suspended in anhydrous toluene, and then a few drops of the mixture were pipetted onto a carbon support film.

.

AI

oC

~,

O AI4C 3 o

3. Results

3.1. XRD analysis The structural evolution as a function of milling time was followed by XRD. Fig. i shows an example for the composition of AIs71C4z.9. After 40 h of milling, the profile lines of At were broadened due to reduction of crystallite size and accumulation of lattice strain. The C peaks disappeared completely, Traces of AI4C~ were observed. After ,o~ h of milling, the intensities oi AI4C 3 peaks remarkably increased, and the intensities of At peaks became weak. On further milling, the Al peaks vanished. After a milling time of 150 h, a broad diffraction maximum, corresponding to an amorphous phase, superimposed by the Bragg peaks of AIaC 3 were found in the XRD pattern. After 220 h milling, only a halo peak corresponding to an amorphous phase was observed. For AI7ECz8, a similar structural evolution process was detected during the first 100 h of milling (Fig. 2). After 150 h of milling, the A1 profile lines besides the A14C3

.6

;r o

11.

220h

20 , AI 0

220h

20

oC

~

• AI4C3

..~-~.~.

40

60 20 (degree)

80

Fig. I. The XRD patterns for A157~C429 powder mixture "after selected times of milling.

A

40

60 2 0 (degree)

80

Fig. 2. The XRD patterns for AIT:C:, after differen! milling times.

peaks remained in XRD patterns. Upon further milling, the material was fully transformed into a single amorphous phase. A different structural evolution process was found during mechanical alloying of the A I 8 4 C ! 6 powder mixture (Fig. 3). After 100 h of milling, AI4C 3 peaks were detected in the XRD pattern. However, the Ai4C 3 peaks gradually disappeared with prolonged milling. After 220 h of milling, only the AI peaks were observed. Assuming that the contributions of the grain size and the strain to the integrated intensities of peaks follow Cauchy and Gaussian distribution respectively, the grain size and the internal strain are then estimated from Cauchy and Gaussian integral breadth component of Voigt function [16]. By applying this method to the A184C10 sample milled for 220 h, the average grain size of Al was calculated to be around 16 nm, and the strain was about 1.30. In addition, the values of aluminum lattice spacing can be estimated from the shift in the diffraction peaks. The lattice parameter of AI for the AIs4C,6 sample milled for 220 h was determined to be 0.4065 rim, which was slightly larger than that of the unmilled elemental A1 (a=0.4049 nm). This can be ascribed to the fact that some carbon atoms were located at interstitial sites in the lattice of Al.

N.Q. Ib'u et al. I Journal oJ Alloys and Compounds 260 (i.o97) 1 2 ! - 1 2 6

~

123

*,AI

r~ °0,,4

,'mq °,.q

¢D

.¢,,o

28 42.9

._L____A 20

40

60 2 6 (degree)

80

Fig. 3. The XRD patterns for AI~4Ct6 after seletted times of milling.

and most of ,hem segregaled in the grain boundaries. This notion is :,upported by the following fac~: The nanocrystalline c~-Al provides a large volume fraction of grain boun,Jaries. And the atomic density in the grain boundaries of Al is appreciably lower than t&at in conventional grain boundaries. In addition, the radius of a C atom is much smaller than ~at of A1 atom (R~=0.0914 nm, RA~=0.143 nm). Therefore it is probable that an appreciable part of carbon atoms are located at the grain boundaries. This is what has also been found in the mechanically alloyed Ti-C alloy [171 and Cr-Fe alloy [18]. Fig. 4 shows the XRD patterns of the five aluminumrich compositioi~s after 220 h of milling. At the carbon concentration up to 23 at%, a single f.c.c, cx-Al was observed. Whereas for compositions 28 at% to 50 at% C, amorphization was found. 3.Z TEM observation

The microstructure of the A157.1C42 9 sample during MA was investigated by TEM (as shown in Fig 5). it must be pointed out that these observations, in some cases, represent the edge of the powder particles due to limitations on the thickness of TEM specimens. As shown in Fig. 5a, the smooth rings of f.c.c. AI are visible in the selected area diffracfio.a pattern (SADP) of the sample milled for 4~) h. The related dark-field nficrograph (Fig. 5a) shows ~:ains with size of less than 18 rim. With prolonged mitring, the amount of AI4C 3 increased. The SADP of the sample milled for 100 h (Fig. 5b) gives nearly continuous rings of AI4C 3 and AI, indicating the random orientation of neighboring crystals. An abundance of nanocrystaUine A l a C 3

j..-~~

._..2___.-0 20

40

6"0 20 (degree)

80

Fig. 4. The XRD patterns for Al~oo_,C , (x=16, 23, 28, 42.9, 50 at%) after 220 h of milling.

grains with a size of about 4-10 nm can be seen in the dark-field image of the (012) reflection (Fig, 5b). Upon further milling, carbide AI4C 3 became amorphous. The SADP of the sample milled for 150 h (Fig. 5c) shows a broad diffraction halo together with the very weak rings of ALC 3. A small number of crystallites are imaged in the amorphous matrix in the corresponding dark-field. In the TEM micrographs t ~ e n after 220 h of milling, the typical feature of an amorphous phase is visible and no crystallites are detected anywhere in the specimen (Fig. 5d), thus confirming the XRD results.

4. Discussion In order to evaluate the driving force for the formation of amorphous phase, thermodynamic calculations have been carried out. Although the stable or metastable equilibria might not be obtained during MA, or the free energy might be raised by the accumulation of defects, the knowledge of free energy is helpful to understand the phase formation during MA. Miedema's model has been applied to estimate the formation enthalpy of the ~_anorphous phase, in Miedema's model the enthalpy of mixing of the solid solution is given by [19,20] z~r-i~ = ~,-/c + M-F ~+

A H s"~'

(1)

where Z~ff/¢. ~ e ~ and z~/S'ruc' refer to chemical, elastic and structural conwibudons to the enthalpy of solid

124

N.Q. Wu et al. I Journal of Alloys and Compounds 260 (1997) 121-126

(d)

Fig. 5. The TEM micrographs for the Al57IC~2~ sample taken at selected milling times. (a) 40 h, (b) 100 h, (c) '='l.,vh, ~uj'"',~,~,.,"""h.

N.Q. Wu et al. / Journal of Alloys and Compounds 260 (1997) 121-126

solution, respectively. The chemical contribution can be obtained by the following equation: ~ei

AH~ = CACa(C~M-IA i, B

+ a,-,s A M S O l

"-'A~'" a in A)

(2)

where ~AI4 . ''j A m a iS the solution enthalpy of A in B, CA and C a are molar fractions of A and B components respectively, C~ is the degree to which A atoms are in contact with B atoms, which is expressed as:

C.V,~,, C~ = | - - C A = ¢ v213 + r v 2'3 ""A-- A

"B-

(3)

B

In the case of the amorphous phase, the elastic and structural contribution to its enthalpy are absent. The total formation enthalpy for the amorphous phase, AH ~, can be estimated as [ 19,20]: AH a = AH¢ + 3.5(CATm, A + CaT,. a)

(4)

where A//~ is obtained by multiplying Eq. (1) by the factor s s 2 [1 +5(CAC B) ], to take account of the presence of chemical short-range order into account [21]. The CALPHAD method has been used to describe the enthalpy of crystalline solid solution phase, and the data required for calculation have been taken from the literature [22-24]. In addition, the enthaply ef AI,tC 3 is given by Ref. [25]. Based on the above described equations, the enthalpycomposition diagr~,n colresponding to 500 K is calculated and plotted in Fig. 6. The temperature of 500 K is an estimation of the effective temperature during the ball milling [26]. In the present work, mechanical alloying of the AI-C

10 A ©

& -1o f.c.c.

-20 m

~0

Ai4C3

40 O

6.2

0.4

0.6

0.8

t

C concentration (.~ole fraction) Fig. 6. The enthalpy-composition diagram for the AI-C system at T=500 K, which is based on ale Miedema's model for the amorphous phase and the CALPHAD method for the f c.c. solid solution.

125

powder mixture fir~Mly results in a supersaturated f.c.c. AI(C) solid solution with a carbon content up to 23 at%, whereas an amorphous phase is formed in the composition range of 28-50 at% C (Fig. 4). This is in agreement with the calculated thermodynamic results on the AI-C system (Fig. 6), where the amorphous phase has lower free energy in the composition range of 27-72 at% C as compared to the f.c.c, solid solution. It should be noticed that the carbide AI4C.~ is preferably fomted before formation of the amorphous phase or the f.c.c, solution during ball milling. The alloying route can be summarized as follows: For A157.1C42.9 For A172C28 For AIa4Ct~

AI + C-->AI4C 3--->amorphous (am.) AI + C--->AI + AIaC 3 AI + am. -->am. AI + C--->AI + AI4C 3-->f.c.c. AI(C)

AI4C 3, having the lowest free energy, preferentially forms in comparison with the other phases. This indicates that the negative free energy of mixing not only provides a thermodynamic driving force for the reaction but also favors the reaction kinetics. In addition, the results demonstrate that the formation of metastable phase during MA originates from destabilization of the compound A14C 3. Therefore, the mechanism of phase selection deviates from metastable phase formation in the thin-film diffusion couple, where the formation of the metastable phase, in particular amorphous phase, is kinetically favorable [11,12]. Once AI,,C 3 is formed, ball milling brings about a large number of defects (such as grain boundaries) in the crystal. It is known that the defects induced by ball milling can increase the free energy of crystal. For example, Fecht et al. have found that the excess enthalpy of about 30 percent of the heat of fusion is stored in the nanocrystalline b.c.c. (or h.c.p.) metals after high-energy ball milling [27]. If the free energy of crystal increases beyond the amorphous phase, the crystal-to-amorphous transition can ~ c u r [28,29]. Jang and Koch [28] have observed that the nanocrystalline structure with the cell diameter of about 2 nm is formed, in some areas of materials, an amorphous structure is detected. In addition, Veprek et al. [30] have reported that an abrupt transition from microcrystalline to amorphous silicon occurs if the size of grain in the chemical vapor deposited silicon film is refined to 3 nm. Based on the experimental results, these authors [28,30] h~ve proposed that the crystalline-to anaorphous transition is induced by increase in grain-boundary energy. As shown in Fig. 6, in At4C3 the difference between the amorphous phase and the crystalline compound is estimated to be 2.6 M/g-atom. If the stored enthaipy in A14C3 exceeds this value, destabilization of A14C3 can occur. Our experiment shows that file nanocrystalline material with a grain size of 4 - l 0 nm was formed before amorphization of AI4C 3 (Fig. 5b). By assuming a grain boundary energy of 0.5 J i m 2, the energy comdbudou due to the grain bound-

126

N.Q. Wu et al. I Journal of Alloys and Compounds 260 (199D 121-126

aei is estimated to be 1.1-2.6 kJ/g-atom. This means that refinement of grain can play an important role in the crystalline-tv-amorphous transformation. Another mechanism suggested for the crystalline-toamorphous transformation of AI4C 3 induced by mechanical milling may be pressure-induced amorphization. During ball milling, it is estimated [15] that the average pressure developed across the contacting surface is in the range of 109 Pa, and the impact time, involving both the ball compression and its subsequent relaxation, is on the order of 10 -5 s [32]. At the relatively rapid unloading rates, the high pressure crystalline phase transforms to a metastable amorphous phase. In support of this notion is the fact that single-crystal silicon and germanium are converted to an amorphous state at room temperature under both Vickers and Knoop indentation [32]. In the indentation experiment, the sharp indenter is loaded onto a flat surface of Si and then unloaded rapidly. It is estimated that the pressure during indentation is on the order of 109 Pa, and the unloading occurs in 30 ms. Clark et al. [33] have proposed that at this rapid unloading that the high pres,,ure, crystalline form of Si cannot transform back fast enough to the equilibrium structure, such that the amorphous phase forms metastably. Comparing the indentation experiment with the ball milling, it is found that the fundamental actions of both the processes are very similar. Therefore, it is not surprising that pressure is responsible for the amorphization of AI.,C~ by mechanical milling. A further supporting experiment indicates that the pressureinduced amorphization has been observed during mechanical milling of polycrystalline Si [31].

5. Conclusions 1. The carbide AI4C~ is firstly ibrmed as an intermediate product during ball milling of AI-C powder blends. Further milling leads to destabilization of AI4C 3. It is proposed that grain-refinement and pressure-induced amorphization are responsible for the crystalline-toamorphous transformation. 2. Bailing milling of an elemental AI-C powder mixture finally results in the f.c.c. AI (C) solid solution with a carbon content up to 23 at%, whereas an amorphous phase is formed in the composition range of 28-50 at%C. 3. The phase ,,;election during MA can be interpreted by the enthalpy-composition diagram.

Acknowledgements This work is finacially supported by the Foundation of Zhejiang University. P.R. China.

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