Thermo-dynamic analysis on solid-state reduction of CaO particles dispersed in Mg–Al alloy

Materials Chemistry and Physics 129 (2011) 631–640

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Thermo-dynamic analysis on solid-state reduction of CaO particles dispersed in Mg–Al alloy Katsuyoshi Kondoh a,∗ , Junji Fujita a , Junko Umeda a , Hisashi Imai a , Keitaro Enami b , Masaki Ohara b , Takanori Igarashi b a b

JWRI, Osaka University, 11-1 Mihogaoka, Ibaragi, Osaka 567-0047, Japan TOPY Industries, Limited, 1 Akemi-cho, Toyohashi-City, Aichi 441-8510, Japan

a r t i c l e

i n f o

Article history: Received 20 August 2010 Received in revised form 15 March 2011 Accepted 8 May 2011 Keywords: Powder metallurgy Heat treatment X-ray diffraction topography Microstructure

a b s t r a c t AZ61B alloy powder composite with Al2 Ca fine dispersoids was developed by using CaO additive particles as raw materials. For synthesis of Al2 Ca in the matrix, CaO particles were elementally and uniformly mixed with AZ61B alloy chips via ECABMA process. The mechanism in formation of Al2 Ca intermetallics was investigated by XRD and SEM–EDS analysis when AZ61B alloy green compacts containing CaO additives were heat treated at 380–625 ◦ C in argon gas atmosphere. A change in a standard free energy in formation of Al2 Ca via reaction between CaO and Mg–Al alloy was calculated by using a standard Gibbs free energy of each element contained in the green compact. Both of the theoretical analysis by thermo-dynamic and experimental investigation clarified that Al2 Ca and MgO were synthesized by employing Mg–Al alloy, not pure Mg as the matrix material. Microstructural analysis indicated that needle-like intermetallic of (Mg,Al)2 Ca or Al3 Ca4 Mg were formed as intermediately created compounds in the solid-state reaction between CaO particles and Mg–Al alloy to synthesize Al2 Ca and MgO dispersoids. © 2011 Elsevier B.V. All rights reserved.

1. Introduction A magnesium (Mg) alloy promises a remarkable lightweight effect due to its low density of about 1.74 g cm−3 , and results in a fuel consumption improvement by its application to structural components of automotives [1–4]. In particular, the heat resistance improvement of Mg alloys is strongly required because a weight reduction of engine blocks and transmission cases used at elevated temperature (120–200 ◦ C) is very important [5–9]. The addition of lanthanum (La), one of rare earth elements, is effective to drastically improve the heat resistance of Mg–Al alloys due to both the formation of Al–La intermetallics (Al11 La3 , Al4 La and Al2 La) [10–15] and decrease of Mg17 Al12 (␤ phase) compounds, which are thermally unstable at the above elevated temperature. As a result, Mg–Al–La cast alloys indicated an excellent creep performance at elevated temperature, which was equivalent to that of the conventional ADC12 aluminum die-cast alloy [14]. However, the application of these alloys to automotive components was very limited because La element was very expensive. The previous studies reported that the calcium (Ca) addition into AZ91 cast alloy had an important benefit in improving the creep resistance at 200 ◦ C by formation of Al2 Ca or (Mg,Al)2 Ca intermetallic compounds existing

∗ Corresponding author. E-mail address: [email protected] (K. Kondoh). 0254-0584/$ – see front matter © 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.matchemphys.2011.05.017

at ␣-Mg grain boundaries [9,16]. This is because these networkstructured Al–Ca intermetallics are useful for improvement of creep performance by preventing the deformation of ␣-Mg grains and grain boundary sliding at elevated temperature. However, a lot of Al–Ca fine particles distributed uniformly in the matrix are much effective in increasing the yield stress of Mg alloys by their pinning effects. When metal Ca chunks are employed as raw materials and melted into molten the Mg–Al alloy, the network-structured Al2 Ca intermetallics are generally formed and distributed at the grain boundaries as mentioned above. It means difficult to disperse Al2 Ca compounds as fine reinforcement particles in the matrix of Mg alloys in use of raw metal Ca materials via casting process. In the present study, calcium oxide (CaO) particles are used as raw materials to form fine Al2 Ca dispersoids in ␣-Mg matrix via solid-state reaction between their particles and the Mg–Al alloy. It is expected that the conventional Mg–Al alloy reinforced with Al2 Ca fine dispersoids shows a high yield stress due to their pinning and also excellent in the cost-performance because of cheap CaO particles used as raw materials. However, from a viewpoint of thermo-dynamic, an Ellingham diagram in the oxide formation [17] indicates that the reduction of CaO by Mg in solid-state (∼650 ◦ C) never occurs and results in no formation of Al2 Ca and MgO. It means the CaO raw particles remain in the cast Mg material. In this preliminary experiment, it was found the above reaction to synthesize Al2 Ca intermetallic compounds occurred by using Mg–6 mass% Al

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K. Kondoh et al. / Materials Chemistry and Physics 129 (2011) 631–640

Fig. 1. Schematic illustration of Equal Channel Angular Balk Mechanical Alloying (ECABMA) process used in fabrication of Mg alloy powder compact containing CaO particles.

alloy (AZ61B) instead of pure Mg. Therefore, first of all, microstructural changes of AZ61B green compacts containing CaO additives after heat treatment are investigated by XRD and SEM–EDS analysis. The thermo-dynamic approach using a change in the standard free energy was applied to clarify the mechanism in solid-state resolution of CaO particles dispersed in the AZ61B green compact during the annealing process at high temperatures.

2. Experimental procedure 2.1. Preparation of AZ61B green compacts with CaO particles AZ61B alloy chips, machined from their original cast ingot (Al; 6.41, Zn; 1.02, Mn; 0.28, Si; 0.02, Fe; 0.004, Cu; 0.002, Ni; 0.0007, Mg; Bal./mass%), were used as raw materials, and had a mean particle size of 1.38 mm measured by particle size analyzer (HORIBA, LA-950). Another input material was CaO particles, having 98.3% purity and a mean particle size of 2.3 ␮m. The elemental mixture of AZ61B chips and CaO particles (0, 2.5, 5.0, 7.5 and 10 vol.%) was mixed by the rocking mill equipment (Seiwa Giken Co., RM-05S) and used as the starting materials. For preparation of AZ61B green compacts containing CaO particles, Equal Channel Angular Balk Mechanical Alloying (ECABMA) process [18], where cold compaction and extrusion are alternately carried out in the die as schematically illustrated in Fig. 1, was employed in the present study. The maximum pressure of 611 MPa was applied in compaction, and the total number of cycles in repeating the compaction and extrusion was 50. ECABMA process was effective to form AZ61B green compacts with uniformly distributed CaO particles in the matrix by the above alternate severe plastic deformation. It had an advantage to break the surface oxide films (MgO) of AZ61B chips and bring newly created Mg surface of the chips into contact with CaO particles. After this process, the columnar green compact (precursor) with 35 mm diameter and about 80 mm length was obtained.

3. Results and discussion 3.1. DTA and XRD analysis results of precursors Fig. 2 shows DTA profiles of the precursors with various contents of CaO particles. When the CaO content was 5 vol.% or more, a small exothermic was obtained at 480 ◦ C and its heat gradually increased with increase in the CaO content. According to XRD analysis results of the precursors with CaO particles shown later, this exothermic heat was caused by a reaction between AZ61 alloy and CaO particles to synthesize Al2 Ca intermetallic compounds. Furthermore, these three precursors containing 5 vol.% or more CaO additives obviously indicated the endothermic behavior at 510 ◦ C. This heat also gradually increased with increasing the CaO content of the precursor. Al–Ca binary phase diagram [19] suggests that this endothermic is due to meltdown of Al2 Ca intermetallic. On the other hand, the precursors with CaO less than 2.5 vol.% showed no exothermic and endothermic mentioned above. According to Al–Ca phase diagram, no Al2 Ca can be formed by the solid-state reaction between CaO and aluminum (Al) elements contained in AZ61B alloy when 2.5 vol.% CaO particles was added in the AZ61B green compact. Fig. 3 indicates XRD patterns of the precursors after annealing, which include 0% (a), 2.5 vol.% (b), 5 vol.% (c), 7.5 vol.% (d) and 10 vol.% (e) CaO particles. As shown in Fig. 3(a), Mg and Mg17 Al12 (␤ phase) peaks, corresponding to the alloy elements of AZ61B raw chips, were detected in the precursor before heat treatment. The precursors annealed at 500 ◦ C and 550 ◦ C revealed a drastic decrease of Mg17 Al12 peaks (䊉) intensity and a little peak shift of

2.2. Heat treatment of green compacts

0.5 0

DTA/m μV.mg-1

For fixing the heat treatment temperature, a differential thermal analysis (DTA, Shimadzu, DTG-60) was conducted for the precursors with various contents of CaO particles under argon gas atmosphere. Argon gas flow rate was 7.5 ml s−1 and heating rate was 5 ◦ C min−1 . As the results by DTA, the annealing treatment from 380 ◦ C to 625 ◦ C in argon gas atmosphere was applied to the precursors by using the tube type furnace for investigation of a reaction behavior between AZ61B matrix and CaO particles. The holding time at 380–500 ◦ C was 3.6 ks, and that at 525–625 ◦ C was 14.4 ks.

-0.5 -1.0 AZ61B

-1.5

AZ61B+2.5vol%CaO

2.3. Microstructural analysis of heat treated compacts

-2.0

AZ61B+5vol%CaO

Optical microscope (OLYNPUS, BX-51P) and field emission scanning electron microscope (FE-SEM, JEOL, JSM-6500F) equipped with energy dispersive X-ray spectrometer (EDS, JEOL, EX-64175 JMU) observation were carried out to investigate microstructural changes of each green compact after annealing. X-ray diffraction analysis (XRD, Shimadzu, XRD-6100) was used for peak identification of the original dispersoids and synthesized compounds of the precursors. Electron probe microanalyzer (EPMA, JEOL, JXA-8600) was used to investigate chemical compositions of the compounds distributed in the matrix of the precursor.

-2.5

AZ61B+10vol%CaO

AZ61B+7.5vol%CaO 5vol%CaO

-3.0 0

100

200

300

400

500

600

700

800

Temperature, T / Fig. 2. DTA profiles of AZ61B powder precursors with CaO particles via ECABMA process.

K. Kondoh et al. / Materials Chemistry and Physics 129 (2011) 631–640

c

Intensity /a.u.

Mg Mg17Al12 MgO Al2Ca CaO Un-known 550

625

Intensity /a.u.

a

633

(1h)

550

(1h) (1h)

500 480 460 440

500

420 400 380

Precursor

30

35

40

45

50

55

Precursor

60

30

35

40

Diffraction angle, 2θ /°

45

50

55

60

Diffraction angle, 2θ /°

b

d

550

(1h) (1h)

500 480 460 440

625

Intensity /a.u.

Intensity /a.u.

625

550 500 480 460 440

420

420

400

400

380

380

Precursor

30

35

40

45

50

55

Precursor

60

30

35

40

Diffraction angle, 2θ /°

45

50

55

60

Diffraction angle, 2θ /°

e Intensity /a.u.

625 550

(1h) (1h)

500 480 460 440 420 400 380 Precursor

30

35

40

45

50

55

60

Diffraction angle, 2θ /° Fig. 3. (a) XRD patterns of AZ61B powder precursors with no CaO particle and after annealing at 500 ◦ C and 550 ◦ C. (b) XRD patterns of AZ61B powder precursors with 2.5 vol.% CaO particles and after annealing at 380–625 ◦ C. (c) XRD patterns of AZ61B powder precursors with 5 vol.% CaO particles and after annealing at 380–625 ◦ C. (d) XRD patterns of AZ61B powder precursors with 7.5 vol.% CaO particles and after annealing at 380–625 ◦ C. (e) XRD patterns of AZ61B powder precursors with 10 vol.% CaO particles and after annealing at 380–625 ◦ C.

them to lower diffraction angle, compared to the as-compacted precursor without heat treatment. Mg–Al binary phase diagram [20] suggests that Mg17 Al12 compounds were resolved by annealing at 500–550 ◦ C and a lot of Al atoms were solid-soluted in ␣-Mg matrix. However, a few of these compounds precipitated during cooling from 500 to 550 ◦ C to room temperature. In addition, MgO peaks were hardly detected in the compacts annealed at 500 ◦ C and 550 ◦ C. It means no remarkable oxidation of Mg elements occurred by annealing less than 550 ◦ C for 3.6 ks under the controlled argon gas atmosphere in the tube furnace. XRD profiles of the precursor with 2.5 vol.% CaO particles shown in Fig. 3(b) indicate no change in CaO peaks () by annealing until 500 ◦ C. However, when the heat treatment at 550 ◦ C or more was employed, CaO peaks disappeared and Al2 Ca and MgO peaks were obviously detected. According to the above results in Fig. 3(a), MgO formation was due to the reaction

between Mg–Al alloy and CaO, not oxidation of Mg during annealing. As shown in Fig. 3(c)–(e), all specimens never showed the peaks of Mg17 Al12 compounds after heat treatment over 380 ◦ C. They also revealed a formation of Al2 Ca and MgO by annealing at 420–500 ◦ C, and their peak intensities gradually increased with increasing the CaO content. However, Al2 Ca peak at 2 = 31.8◦ disappeared by annealing at 550 ◦ C or more. DTA profiles in Fig. 2 indicate the peak endothermic temperature due to the meltdown of Al2 Ca compound is about 545 ◦ C, which corresponds to its eutectic point in Al–Ca binary phase diagram as shown in Fig. 4. Therefore, a meltdown of Al2 Ca intermetallics was caused by heat treatment at 550 ◦ C or more while Al2 Ca compounds were synthesized via solid-state reaction between AZ61B and CaO during annealing at 420–500 ◦ C, and resulted in the peak identification of MgO, not Al2 Ca. The heat treatment at 550 ◦ C and 625 ◦ C applied to the precursors caused

634

K. Kondoh et al. / Materials Chemistry and Physics 129 (2011) 631–640

0.05

Intensity ratio Al2Ca / Mg

Annealing time; 14.4 ks 10vol% 10vol %

0. 0 04

7.5vol%

0.03 5vol%

0. 02 0.01 0 360

2.5vol%

380

400

420

440

460

480

500

520

Annealing temperature, T / Fig. 5. Dependence of peak intensity ratio of Al2 Ca and Mg peaks on annealing temperature of AZ61B powder precursors with various contents of CaO particles.

Fig. 4. Al–Ca binary equilibrium phase diagram.

a disappearance of both CaO and Al2 Ca peaks, but Mg, MgO and un-known peaks were detected. In particular, the intensity of MgO peak (2 = 43.3◦ ) of these precursors became much stronger than that of specimens with heat treatment less than 500 ◦ C. This is because solid-soluted oxygen atoms reacted with Mg elements and resulted in an acceleration of MgO formation. In addition, some of un-known peaks () were obviously detected in the precursor annealed at 550 ◦ C and 625 ◦ C. The identification of chemical compositions of these un-known phases was discussed in EPMA results later. Furthermore, as shown in Fig. 3(b), Al2 Ca peaks were clearly detected in the precursors with 2.5 vol.% CaO particles when the heat treatment at 550 ◦ C and 625 ◦ C was employed. The DTA profile of this specimen in Fig. 2 clarified that a meltdown of ␣-Mg occurred at about 530 ◦ C Accordingly, Al2 Ca compounds detected in the precursor with 2.5 vol.% CaO annealed at 550–625 ◦ C were synthesized by the liquid-state reaction between AZ61B and CaO.

This mechanism in Al2 Ca formation was quite different from that in the case of the precursors containing 5 vol.% or more CaO particles. Fig. 4 suggests that Al2 Ca compounds are melt and disappear by heat treatment at 700 ◦ C or more when the content of CaO particles dispersed in AZ61B compact is 2.5 vol.%. Therefore, Al2 Ca and MgO compounds are still remain in the matrix of the precursors after annealing at 550 ◦ C and 625 ◦ C for 3.6 ks. According to the above XRD profiles, the quantitative evaluation of Al2 Ca formation during heat treatment was conducted. The integrated peak intensity of ␣-Mg (2 = 36.6◦ ) and Al2 Ca (2 = 31.8◦ ) were calculated by XRD analysis results shown in Fig. 3, and a dependence of their peak intensity ratio (Al2 Ca/Mg) on the heat treatment temperature is estimated in Fig. 5. As mentioned above, the reaction behavior between AZ61B and CaO at 550 ◦ C or more was caused by the content of CaO additives of the precursors. Accordingly, the formation analysis of Al2 Ca during heat treatment at 380–500 ◦ C was investigated. As shown in Fig. 5, all precursors excepting 2.5 vol.% CaO content revealed the increase of in situ formed Al2 Ca compounds via solid-state reaction during annealing with increase in the temperature. The peak intensity ratio of

Fig. 6. SEM observation of AZ61B powder precursors with various contents of CaO additive particles before heat treatment. CaO content of 2.5 vol.% (a), 5 vol.% (b), 7.5 vol.% (c) and 10 vol.% (d).

K. Kondoh et al. / Materials Chemistry and Physics 129 (2011) 631–640

635

Fig. 7. SEM observation of AZ61B powder precursors with various contents of CaO additive particles after annealing at 420 ◦ C for 14.4 ks. CaO content of 2.5 vol.% (a), 5 vol.% (b), 7.5 vol.% (c) and 10 vol.% (d).

Fig. 8. SEM–EDS analysis result of AZ61B powder compact with 10 vol.% CaO particles after annealing treatment at 500 ◦ C for 14.4 ks.

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K. Kondoh et al. / Materials Chemistry and Physics 129 (2011) 631–640

Fig. 9. SEM–EDS analysis result of AZ61B powder compact with 10 vol.% CaO particles after annealing treatment at 625 ◦ C for 3.6 ks.

3.2. Microstructural analysis results Fig. 6 shows SEM observation of AZ61B green compacts with various contents of CaO particles before heat treatment. CaO particles, having a particle size of about 150 nm to 4 ␮m in diameter, were uniformly dispersed in the matrix. They mean that a remarkable fragmentation of CaO raw particles occurred by the repetition of severe plastic deformation in ECABMA process. Mg17 Al12 compounds were also mechanically broken into small dispersoids with 1–5 ␮m via ECABMA process. Fig. 7 indicates SEM photos of the precursors annealed at 420 ◦ C for 14.4 ks. As shown in (a), raw materials of CaO and Mg17 Al12 intermetallics were observed, and no formation and precipitation of Al2 Ca compounds was detected in the matrix. This result corresponds to the XRD profile shown in Fig. 3(b). In the case of CaO content with 5 vol.% or more, Al2 Ca compounds were obviously observed, and their number gradually increased with increasing the CaO content. However, a few CaO additive particles still remained to be not reacted with Al in

the matrix. This means the incomplete reaction between CaO and AZ61B occurred of each precursor during annealing at 420 ◦ C, and also agrees with XRD results in Fig. 3(c)–(e). Fig. 8 shows SEM–EDS analysis results of AZ61B + 10 vol.% CaO compact after annealing at 500 ◦ C for 14.4 ks. The distribution area of Al elements correStandard Gibbs free energy, ΔG0 (kJ/mol)

AZ61B–10 vol.% CaO compact was saturated with heat treatment at 420 ◦ C or more, and that of the precursor containing 7.5 vol.% CaO indicated its saturation by annealing at 480–500 ◦ C. These results mean that Ca elements originated in CaO particles were completely reacted with Al elements of AZ61B when its ratio reached a saturated constant in each specimen.

0

Ellingham diagram -200

Calculated value

-400

2Mg+O2 = 2MgO -600

4/3Al+O2 = 2/3Al2O3 -800 -1000

2Ca+O2 = 2CaO

-1200 -1400 0

500

1000

1500

2000

2500

Temperature, T / Fig. 10. Changes in standard Gibbs free energy (G◦ ) of Al2 O3 , MgO and CaO formation by using thermo-dynamic database compared to Ellingham diagram in each oxide formation.

Standard Gibbs free energy, ΔG0 (kJ/mol)

K. Kondoh et al. / Materials Chemistry and Physics 129 (2011) 631–640 Table 1 Ca/Al atomic ratio for each Al–Ca–Mg intermetallic compound.

100

(1) Mg + CaO = Ca + MgO

50 0

(2) 3Mg + CaO = Mg2Ca + MgO -50

Atomic ratio

Ca/Al

Al2 Ca Al4 Ca Al14 Ca13 Al3 Ca8 (Mg,Al)2 Ca Ca4 Al3 Mg

0.25 0.25 0.93 2.67 0.25 1.33

-100

(3) Mg + 2Al + CaO = Al2Ca + MgO -150 -200 0

200

400

600

800

1000 1200 1400

1600

Temperature, T / Fig. 11. Changes in standard Gibbs free energy (G◦ ) of each chemical reaction in Mg–Al–CaO system by using thermo-dynamic database.

Intensity /a.u.

(e) 14.4 ks

Mg Mg17Al12 MgO Al2Ca CaO Un

(d) 3.6 ks (c) 600 s

(a) 0 s

35

40

45

50

sponded well to that of Ca elements. It suggests in situ formed Al2 Ca fine compounds less than 2 ␮m were uniformly dispersed in the matrix. As mentioned in the XRD profiles of Fig. 3(c)–(e), un-known peaks were observed in the precursors via heat treatment at 550 ◦ C and 625 ◦ C as shown in Fig. 9. SEM–EDS analysis was carried out on the AZ61B + 10 vol.% CaO green compact after annealing at 625 ◦ C for 3.6 ks. A typical microstructure of this specimen consisted of fine particles (Area A) and lamellar-structured intermetallics. By considering the EDS mapping and XRD profile of Fig. 3(e), the former particles correspond to MgO, and the lamellar compounds are detected as Al–Ca–Mg intermetallics containing no oxygen. In addition, EPMA investigation applied to 10 points of lamellar intermetallics indicated the atomic ratio of Ca and Al (Ca/Al) was 1.25–2.22. According to Table 1 showing the Ca/Al values of Al–Ca and Al–Ca–Mg intermetallics reported in the previous studies [21,22], the lamellar-structured intermetallics distributed in the matrix shown in Fig. 9 are detected as (Mg,Al)2 Ca intermetallic. 3.3. Reaction behavior of CaO and AZ61B during annealing

(b) 300 s

30

637

55

60

Diffraction angle, 2θ /° Fig. 12. XRD patterns of AZ61B powder precursors including 10 vol.% CaO particles after annealing treatment at 500 ◦ C for 0 s, 300 s, 600 s, 3.6 ks and 14.4 ks.

The above investigation results suggest CaO additive particles were reduced by Mg in solid-state and Al2 Ca and MgO were synthesized when the CaO content was 5 vol.% or more. According to Ellingham diagram, the standard Gibbs free energy (G) in formation of calcium oxide (CaO) at 0–2000 ◦ C is much smaller than those in magnesium oxide (MgO) and aluminum oxide (Al2 O3 ) formation. This means no reduction of CaO by Mg at 400–500 ◦ C (solid-state),

Fig. 13. SEM microstructure observation of AZ61B powder precursors including 10 vol.% CaO particles after annealing treatment at 500 ◦ C for 0 s (a), 300 s (b), 600 s (c) and 3.6 ks (d).

638

K. Kondoh et al. / Materials Chemistry and Physics 129 (2011) 631–640

Table 2 Standard Gibbs free energy of various elements, oxides and intermetallics referenced from thermo-dynamic database used in this study.

T/°C

Mg

Al

Ca

O2

MgO

CaO

Al2O3

Al2Ca

Mg2Ca

25

-9.7

-8.4

-12.4

-61.2

-606.3

-646.5

-1690.9

-245.1

-70.1

27

-9.8

-8.5

-12.5

-61.5

-606.4

-646.6

-1691.0

-245.3

-70.3

127

-13.5

-11.7

-17.0

-82.5

-609.8

-651.1

-1697.4

-255.0

-81.9

227

-17.8

-15.6

-22.3

-104.3

-614.3

-656.7

-1706.3

-266.7

-95.5

327

-22.7

-20.0

-28.1

-126.6

-619.8

-663.4

-1717.4

-280.0

-110.6

427

-28.0

-24.8

-34.4

-149.6

-626.1

-670.9

-1730.4

-294.8

-127.2

527

-33.8

-30.1

-41.3

-172.9

-633.2

-679.2

-1745.0

-310.8

-144.9

627

-39.9

-35.8

-48.7

-196.7

-640.9

-688.1

-1761.1

-327.9

-163.8

727

-47.2

-42.6

-56.4

-220.9

-649.1

-697.6

-1778.6

-346.0

-183.7

827

-54.9

-50.1

-64.6

-245.4

-657.9

-707.6

-1797.4

-365.0

927

-63.0

-57.9

-73.7

-270.3

-667.2

-718.2

-1817.3

-384.8

1027

-71.4

-65.9

-83.3

-295.4

-676.9

-729.2

-1838.2

-405.3

1127

-83.2

-74.2

-93.1

-320.8

-687.0

-740.6

-1860.2

-428.5

1227

-101.4

-82.7

-103.2

-346.5

-697.5

-752.4

-1883.0

-454.3

1327

-119.6

-91.5

-113.5

-372.4

-708.4

-764.6

-1906.8

-480.8

1427

-138.1

-100.4

-124.1

-398.6

-719.6

-777.2

-1931.4

-507.9

1527

-156.6

-109.5

-137.0

-424.9

-731.2

-790.1

-1956.7

-535.5

1627

-175.2

-118.8

-156.3

-451.5

-743.0

-803.3

-1982.8

1727

-194.0

-128.3

-175.7

-478.3

-755.2

-816.8

-2009.7

1827

-137.9

-505.2

-767.7

-830.6

-2037.2

1927

-147.7

-532.4

-780.4

-844.7

-2065.4

2027

-157.6

-559.7

-793.4

-859.0

-2094.2

Black: solid phase; blue: liquid phase: red; gas phase.

and results in no possibility to form Al2 Ca intermetallics from a viewpoint of thermo-dynamic analysis. On the other hand, the thermo-dynamic database [23,24] shown in Table 2 was useful to calculate the standard Gibbs free energy of various inorganic compounds. In the case of the chemical reaction to form materials C and D via reaction between materials A and B (aA + bB → cC + dD), a change in the standard free energy (G◦ ) is calculated by the following equation: ◦ ) − (aGA◦ + bGB◦ ) G◦ = (cGC◦ + dGD

GA◦ ,

GB◦ ,

GC◦

(1)

◦ GD

and are a standard Gibbs free energy of each where material. First of all, for evaluation of the accuracy and reliability of this database, G◦ values of Al2 O3 , MgO and CaO are calculated by using these data in Table 2, and compared to those of Ellingham diagram [17]. As shown in Fig. 10, the estimation values in each oxide formation using the database reveal a good agreement with the traditional Ellingham diagram data. This result means the database has a high reliability for the thermo-dynamic analysis to investigate the reaction behaviors. Raw materials used in the present study suggest the following reaction equations in reduction of CaO by Mg–Al alloy: (i) Mg + CaO = Ca + MgO (ii) 3Mg + CaO = Mg2 Ca + MgO (iii) Mg + 2Al + CaO = Al2 Ca + MgO According to the above reaction equations, G◦ values to reduce CaO in Mg–Al–CaO system were calculated by using the database in Table 2, and plotted as a function of temperature in Fig. 11. In Eqs.

(i) and (ii), G◦ reveals positive values in the temperature range investigated in this study. That is, from a thermo-dynamic point of view, these two reactions never occur. However, G◦ values of equation (iii) are completely negative, and results in the progress of the reduction of CaO particles by Mg–Al alloy. It means that CaO is resolved into Mg–Al alloy and Al2 Ca and MgO are formed when heat treatment is applied to Mg–Al–CaO system. In addition, the calculation result in Fig. 11 also indicates Al2 Ca can be synthesized at room temperature by a reaction expressed as Eq. (i). However, the experimental results showed the suitable annealing temperature at 420–500 ◦ C was necessary for solid-state synthesis of Al2 Ca compounds. For clarification of the difference between them, the diffusion mechanism of Al atoms to CaO particles must be investigated in detail. 3.4. Effect of annealing time on microstructures The effect of the holding time in heat treatment on microstructure and phase changes of the precursor was investigated. The content of CaO additives contained in AZ61B green compacts was 10 vol.% and heat treatment was conducted at 500 ◦ C for 0, 300, 600, 3.6 and 14.4 ks in argon gas atmosphere. Fig. 12 shows XRD profiles of the precursors via heat treatment at 500 ◦ C with a different holding time, and Fig. 13 indicates microstructures of each specimen by SEM observation. In the case of 0 s holding time, very small peaks of Al2 Ca and MgO were detected, but the peaks of CaO raw materials were mainly observed in the XRD pattern. SEM observation also indicates that CaO particles clearly existed in the matrix, but some of Al–Ca small compounds were detected around CaO particles. When the holding time of 300 s and 600 s

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Fig. 14. SEM–EDS analysis result of needle-like intermetallics dispersed in matrix of AZ61B powder precursors including 10 vol.% CaO particles after annealing treatment at 500 ◦ C for 300 s.

was employed, no Al2 Ca but un-known peaks were recognized in the XRD profiles of Fig. 12(b) and (c). SEM photos in Fig. 13(b) and (c) also show a lot of needle-like compounds with 2–5 ␮m length dispersed in the matrix, which correspond to un-known phase materials and never observed in (a). Fig. 14 reveals SEM–EDS analysis result of the needle-like intermetallic contained in the precursor after annealing at 500 ◦ C for 300 s, and means that this dispersoid is Al–Ca or Al–Ca–Mg intermetallic. Furthermore, EPMA result of this intermetallic shows that Ca/Al atomic ratio of the needle-like intermetallic was 0.6–1.5 and its average value of about 0.81. In comparison of the values of Al–Ca–Mg intermetallics shown in Table 1, the needle-like intermetallic was detected as (Mg,Al)2 Ca or Al3 Ca4 Mg [25]. With increase in the holding time of 3.6 ks or more, un-known peaks completely disappeared in XRD profiles of Fig. 12(d) and (e), and instead of them, Al2 Ca and MgO were detected in the XRD and SEM observation results. In addition, no needle-like compound existed in the matrix after annealing at 500 ◦ C for 3.6 ks in Fig. 13(d). According to the above microstructures and phase changes, (Mg,Al)2 Ca or Al3 Ca4 Mg intermetallics with needle-like morphology were formed as intermediately created compounds in the solid-state reaction between Mg–Al alloy with CaO particles to synthesize Al2 Ca and MgO dispersoids. As a result in consideration of the above thermo-dynamic analysis and this microstructure observation, the solid-state synthesis of Al2 Ca

intermetallic compounds occurred as the following reaction: Mg + 2Al + CaO → (Mg, Al)2 CaorAl3 Ca4 Mg + MgO → Al2 Ca + MgO

4. Conclusion AZ61B alloy green compacts containing CaO additive particles via ECABMA process were heat treated to synthesize Al2 Ca intermetallic compounds in solid-state. XRD and SEM–EDS analyses were applied to investigate the effect of Al elements of contained in the matrix on the reduction mechanism of CaO particles. A change in a standard free energy in formation of Al2 Ca via reaction between CaO and Mg–Al alloy was calculated by using a standard Gibbs free energy of each element contained in the green compact. In the use of pure Mg, the reduction of CaO by Mg never occurred from viewpoints of both thermo-dynamic analysis and experimental results. However, when employing Mg–Al alloy instead of pure Mg, the calculation result indicated that the change in a standard free energy became negative, and the experimental result also showed a formation of Al2 Ca and MgO dispersoids in the matrix. Microstructural analysis indicated that needle-like intermetallic of (Mg,Al)2 Ca or Al3 Ca4 Mg were formed as intermediately created compounds in

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