Phase transformations in a Cu-25at.%Al-7.5at.%Mn alloy

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Phase transformations in a Cu–25at.%Al–7.5at.%Mn alloy

S.Y. Yang, T.F. Liu

*

Department of Materials Science and Engineering, National Chiao Tung University, 1001 Ta Hsueh Road, Hsinchu 30049, Taiwan, ROC Received 24 August 2005; received in revised form 30 September 2005; accepted 27 October 2005

Available online 29 November 2005

Abstract

The as-quenched microstructure of the Cu–25at.%Al–7.5at.%Mn alloy was D03phase containing extremely ﬁne L–J precipitates (the

L–J phase is a new type of precipitate, which was ﬁrstly observed and identiﬁed by Liu and Jeng (designated L–J phase) in a Cu2.2

-Mn0.8Al alloy.). When the as-quenched alloy was aged at temperatures ranging from 500C to 700 C, the phase transition sequence

was found to be (c-brass + L–J + D03)! (c-brass + L–J + B2) ! b. This result is diﬀerent from that reported by previous workers

in Cu–25at.%Al–(0–8)at.%Mn alloys.

Keywords: Cu–Al–Mn alloy; D03; L–J phase; c-brass; B2

1. Introduction

In previous studies, it is seen that when an alloy with a chemical composition in the range of Cu–25at.%Al–(0– 8)at.%Mn was solution-treated in single b phase (disor-dered body-centered cubic) region and then quenched rapidly, a b! B2 ! D03phase transition occurred during

quenching[1–3]. When the as-quenched alloys were aged at temperatures ranging from 350C to 600 C, c-brass parti-cles were found to appear within the D03or B2 matrix[4–

6]. The c-brass has a D83 (ordered body-centered cubic)

structure with lattice parameter a = 0.872 nm [7,8], and the orientation relationship between the c-brass and the matrix was cubic to cubic[9,10]. Based on their studies, it is seen that the stable microstructure of the Cu– 25at.%Al–(0–8)at.%Mn alloys as the aging temperature increased changed from (c-brass + D03)! (c-brass + B2)

! b. Recently, we made transmission electron microscopy (TEM) observations of the phase transformations of the Cu–25at.%Al–7.5at.%Mn alloy. Consequently, it was found that in the as-quenched condition, the microstruc-ture of the alloy was D03phase containing extremely ﬁne

L–J precipitates. The L–J phase has an orthorhombic structure with lattice parameters a = 0.413 nm, b = 0.254 nm and c = 0.728 nm [11]. When the as-quenched alloy was aged at temperatures ranging from 500C to 700C, the phase transition sequence was found to be (c-brass + L–J + D03)! (c-brass + L–J + B2) ! b, rather

than the (c-brass + D03)! (c-brass + B2) ! b reported

by previous workers in Cu–25at.%Al–(0–8)at.%Mn alloys. 2. Experimental procedure

The alloy, Cu–25at.%Al–7.5at.%Mn (Cu2.7Mn0.3Al),

was prepared in a vacuum induction furnace by using 99.9% Cu, 99.9% Al and 99.9% Mn. The melt was chill cast into a 30· 50 · 200-mm copper mold. After being homo-genized at 900C for 72 h, the ingot was sectioned into 2.0-mm thick slices. These slices were subsequently solu-tion-treated at 900C for 1 h and then quenched into room-temperature water rapidly. The aging processes were performed at temperatures ranging from 500C to 700 C for various times in a vacuum heat-treated furnace and then quenched rapidly.

TEM specimens were prepared by means of a double-jet electropolisher with an electrolyte of 70% methanol and 30% nitric acid. The polishing temperature was kept in

* _{Corresponding author. Tel.: +886 3 5731675; fax: +886 3 5728504.}

E-mail address:tﬂ[email protected](T.F. Liu).

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the range from30 C to 15 C, and the current density was kept in the range from 3.0· 104

A/m2to 4.0· 104

A/m2. Electron microscopy was performed on a JEOL JEM-2000FX scanning transmission electron microscope operat-ing at 200 kV. This microscope was equipped with a Link ISIS 300 energy-dispersive X-ray spectrometer (EDS) for chemical analysis. Quantitative analyses of elemental con-centrations for Cu, Al and Mn were made with the aid of a Cliﬀ–Lorimer ratio thin section method.

3. Results and discussion

Fig. 1(a) is a bright-field (BF) electron micrograph of the as-quenched alloy, showing that a high density of extremely fine precipitates was formed within the matrix.Fig. 1(b) and (c) are two selected-area diffraction patterns (SADPs) of the as-quenched alloy. It is seen in the SADPs that in addition to the reflection spots corresponding to the D03 phase

[12,13], the diﬀraction patterns also consist of extra spots with streaks caused by the presence of the extremely ﬁne

precipitates. When compared with our previous studies in Cu2.2Mn0.8Al and Cu–14.6Al–4.3Ni alloys [11,14], it is

found that these extra spots with streaks should belong to the L–J phase with two variants.Fig. 1(d) is að1 1 1Þ D03

dark-ﬁeld (DF) electron micrograph of the same area as

Fig. 1(a), revealing the presence of the ﬁne D03 domains

with a/2h1 0 0i anti-phase boundaries (APBs). Fig. 1(e), a (0 0 2) D03DF electron micrograph, shows the presence of

the small B2 domains with a/4h1 1 1i APBs. In Fig. 1(d) and (e), it is seen that the sizes of both D03and B2 domains

are very small. Therefore, it is deduced that the D03phase

existing in the as-quenched alloy was formed by a b! B2 ! D03 continuous ordering transition during

quenching [15,16]. Fig. 1(f) is a (1 0 01) L–J DF electron

micrograph, exhibiting the presence of the extremely ﬁne L–J precipitates. Based on the above observations, it is con-cluded that the as-quenched microstructure of the alloy was D03phase containing extremely ﬁne L–J precipitates, where

the D03phase was formed by the b! B2 ! D03

continu-ous ordering transition during quenching.

Fig. 1. Electron micrographs of the as-quenched alloy: (a) BF, (b) and (c) two SADPs. The zone axes of the D03phase are (b) [1 0 0] and (c) [1 1 0],

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When the as-quenched alloy was aged at 500C for moderate times, some coarse particles with a cubic shape started to occur. A typical example is shown in Fig. 2.

Fig. 2(a) shows a BF electron micrograph of the alloy aged at 500C for 20 min. Electron diﬀraction demonstrated that the cubic-shaped particles were of c-brass phase.

Fig. 2(b) is an SADP taken from an area covering the

particle marked as ‘‘R’’ in Fig. 2(a) and its surrounding matrix. It indicates that the orientation relationship between the c-brass and the D03 matrix is (0 0 1)c-brass

k-(0 0 1)m and (0 1 0)c-brassk(0 1 0)m, which is similar to that

found by previous workers in Cu–Al based alloys [9,10].

Fig. 2(c), að1 1 1Þ D03DF electron micrograph of the same

area asFig. 2(a), demonstrates that the D03domains had

grown considerably and that the a/2h1 0 0i APBs gradually disappeared. InFig. 2(c), it is also seen that the c-brass par-ticles occurred preferentially at a/2h1 0 0i APBs. Shown in

Fig. 2(d) is a (1 0 01) L–J DF electron micrograph, revealing

that after being aged at 500C for 20 min, the size of the L–J precipitates was increased signiﬁcantly. With increas-ing the agincreas-ing time at 500C, the c-brass particles grew rap-idly and their morphology changes from cubic to irregular shape, as illustrated in Fig. 3(a). Fig. 3(b) is an SADP taken from a region marked as ‘‘A’’ inFig. 3(a), indicating that the intensity of the reﬂection spots and streaking behavior of the L–J precipitates increased with increasing aging time.Fig. 3(c), a (1 0 01) L–J DF electron micrograph,

reveals that the L–J precipitates grew considerably and the L–J precipitates surrounding the c-brass particles were much larger than those away from c-brass particles.

Fig. 3(d) is að1 1 1Þ D03DF electron micrograph, showing

that the D03 domains have grown to reach a complete

grain. Apparently, the microstructure of the alloy present at 500C was a mixture of (c-brass + L–J + D03) phases.

It is noted here that the coexistence of the (c-brass + L–J) phases has never been observed by previous workers in the Cu–Al–Mn alloy systems before.

Shown in Fig. 4(a) is a (1 0 01) L–J DF electron

micro-graph of the alloy aged at 600C for 1 h and then quenched, indicating that the coexistence of (c-brass + L– J) phases could also be observed. However, it is clearly seen in Fig. 4(a) that two types of L–J precipitates could be detected: one is the larger L–J precipitates surrounding the c-brass particles which were existent at the aging tem-perature, and the other is the extremely ﬁne L–J precipi-tates (the size being comparable to that observed in the as-quenched alloy) which were formed during quenching from the quenching temperature. Fig. 4(b) and (c) show ð1 1 1Þ and (0 0 2) D03 DF electron micrographs, clearly

exhibiting small quenched-in D03domains and well-grown

B2 domains, respectively. This indicates that the micro-structure of the alloy present at 600C was a mixture of (c-brass + L–J + B2) phases. TEM observations indicated that the (c-brass + L–J + B2) was preserved up to 675C. However, when the alloy was aged at 700C and then quenched, the microstructure is similar to that observed in the as-quenched alloy. This indicates that the micro-structure existing at 700C or above is the single disor-dered b phase.

Based on the preceding results, it is obvious that the as-quenched microstructure of the present alloy was D03

phase containing extremely ﬁne L–J precipitates. This is diﬀerent from that examined by other workers in the Cu–25at.%Al–(0–8)at.%Mn alloys, in which they reported

Fig. 2. Electron micrographs of the alloy aged at 500C for 20 min: (a) BF, (b) an SADP. The zone axis of the D03phase is [1 0 0] (h k l = D03, h k l1= L–J

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that the as-quenched microstructure was single D03phase

[1–3]. Here, it is worthwhile pointing out that the extremely ﬁne L–J precipitates had also been detected and identiﬁed by the present workers in the as-quenched Cu2.2Mn0.8Al

and Cu2MnAl alloys [11,17]. However, compared to our

previous studies, it is clear that the amount of the L–J pre-cipitates existing in the present alloy is less than that observed in the previous alloys. It seems to imply that

the higher Mn content in the Cu3xMnxAl alloys may

enhance the formation of the extremely ﬁne L–J precipi-tates within the matrix during quenching.

The coexistence of (c-brass + L–J) phases is a remark-able feature in the present study, which has never been observed by other workers in the Cu–Al–Mn alloy systems before. In order to clarify this feature, an STEM-EDS study was performed. Fig. 5(a) and (b) represent two

Fig. 3. Electron micrographs of the alloy aged at 500C for 2 h: (a) BF, (b) an SADP. The zone axis of the D03phase is [1 0 0] (h k l = D03, h k l1= L–J

phase and h k l = c-brass), (c) and (d) (1 0 01) L–J andð1 1 1Þ D03DF, respectively.

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typical EDS spectra for a c-brass particle and an L–J pre-cipitate in the alloy aged at 500C for 2 h, respectively. The quantitative analyses revealed that the atomic percentages of the alloying elements in the c-brass particle and L–J precipitate were Cu–29.84at.%Al–2.76at.%Mn and Cu– 17.52at.%Al–15.34at.%Mn. It is clear that the concentra-tion of Mn in the c-brass is much less than that in the as-quenched alloy. Therefore, it is expected that along with the growth of c-brass particles, the surrounding regions would be enriched in Mn. The enrichment in Mn would cause the Mn-rich L–J precipitates to form at the regions contiguous to the c-brass particles, as observed in Figs. 3(c) and 4(a).

4. Conclusions

1. In the as-quenched condition, the microstructure of the Cu–25at.%Al–7.5at.%Mn alloy was D03phase

contain-ing extremely ﬁne L–J precipitates.

2. When the as-quenched alloy was aged at temperatures ranging from 500C to 675 C, c-brass particles were found to occur preferentially at APBs. With increasing aging time, the L–J precipitates started to appear at the regions contiguous to the c-brass particles. The coexistence of (c-brass + L–J) phases has never been observed by other workers in the Cu–Al–Mn alloy systems before.

Acknowledgments

The authors are pleased to acknowledge the ﬁnancial support of this research by the National Science Council, Republic of China under Grant NSC93-2216- E009-016. They are also grateful to M.H. Lin for typing the manuscript.

References

[1] Marcos J, Vives E, Casta´n T. Phys Rev B 2001;63:224418.

[2] Kainuma R, Satoh N, Liu XJ, Ohnuma I, Ishida K. J Alloy Compd 1998;266:191.

[3] Prado M, Sade M, Lovey F. Scripta Metall Mater 1993;28:545. [4] Obrado´ E, Frontera C, Maon˜sa L, Planes A. Phys Rev B

1998;58:14245.

[5] Counioux JJ, Macqueron JL, Robin M, Scarabello JM. Scripta Metall 1988;22:821.

[6] Miettinen J. Calphad 2003;27:103.

[7] Dvorack MA, Kuwano N, Polat S, Chen H, Wayman CM. Scripta Metall 1983;17:1333.

[8] Kozubski R, Soltys J. J Mater Sci 1982;17:1441.

[9] Singh J, Chen H, Wayman CM. Scripta Metall 1985;19:887. [10] Dutkiewicz J, Pons J, Cesari E. Mater Sci Eng 1992;A158:119. [11] Jeng SC, Liu TF. Metall Mater Trans 1995;26A:1353. [12] Kuwano N, Wayman CM. Metall Trans 1984;15A:621. [13] Wu CC, Chou JS, Liu TF. Metall Trans 1991;22A:2265. [14] Tan J, Liu TF. Mater Chem Phys 2001;70:49.

[15] Swann PR, Duﬀ WR, Fisher RM. Metall Trans 1972;3:409. [16] Allen SM, Chan JW. Acta Metall 1976;24:425.

[17] Chu KC, Liu TF. Metall Mater Trans 1999;30A:1705. Fig. 5. Two typical EDS spectra obtained from (a) a c-brass particle, and