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BULETIN APLINDO N0.48/2016, April - Mei 2016
Asosiasi Industri Pengecoran Logam Indonesia
Gedung Manggala Wanabakti Blok IV Lantai 3 Ruang 303A
1. School of Materials Science and Engineering, Tsinghua University, Beijing 100084, China; 2. Key Laboratory for Advanced Materials Processing Technology (Ministry of Education), Beijing 100084, China
Abstract: As an excellent giant-magnetostrictive material, Tb-Dy-Fe alloys (based on Tb0.27-0.30Dy0.73-
0.70Fe1.9-2
Laves compound) can be applied in many engineering fields, such as sonar transducer systems, sensors, and
micro-actuators. However, the cost of the rare earth elements Tb and Dy is too high to be widely applied for the
materials. Nowadays, there are two different ways to substitute for these alloying elements. One is to partially
replace Tb or Dy by cheaper rare earth elements, such as Pr, Nd, Sm and Ho; and the other is to use non-rare
earth elements, such as Co, Al, Mn, Si, Ce, B, Be and C, to substitute Fe to form single MgCu2-type Laves phase
and a certain amount of Re-rich phase, which can reduce the brittleness and improve the corrosion resistance of
the alloy. This paper systemically introduces the development, the fabrication methods and the corresponding
preferred growth directions of Tb-Dy-Fe alloys. In addition, the effects of alloying elements and heat treatment on
magnetostrictive and mechanical properties of Tb-Dy-Fe alloys are also reviewed, respectively. Finally, some
possible applications of Tb-Dy-Fe alloys are presented.
Fig. 7: Ratio (λ///Wh) of magnetostriction to hysteresis for x(Tb0.15Ho0.85Fe2)+(1-x)(Tb0.3Dy0.7Fe2) alloys in
different compositions at a magnetic field of 320 kA·m-1 [44]
Therefore, the addition amount of Pr should not exceed 20%, otherwise it is easy to form impurity phase [49]
.
RenZhi et al. [15]
studied the structure and magnetostriction of PrxTb0.2Dy0.8 -xFe1.85C0.05 (x=0.1-0.4) alloys. The
research shows that RFe3 phase and rare earth phase appeared when x≥0.2, which leads to the decrease of
magnetostriction coefficient and Curie temperature.
Figure 8 depicts the magnetostriction coefficient and Curie temperature of the PrxTb0.2Dy0.8-xFe1.85C0.05 (x=0.1-0.4) alloy, and it can be seen that Pr0.2Tb0.2Dy0.6Fe1.85C0.05 alloy shows good magnetostrictive properties. Adding B into TbDyPrFe alloys can restrain the formation of RFe3, therefore it can increase the amount of Pr to 30%. W. J. Ren et al.
[50] studied the TbxDy0.7-xPr0.3(Fe0.9B0.1)1.93 alloy, and the result showed that Tb0.25Dy0.45Pr0.3
(Fe0.9B0.1)1.93 alloy possesses excellent magnetostrictive properties with λ111≈1,850 ppm. Moreover, W. J. Ren et al.
[51, 52] investigated Tb0.2Dy0.82xPrx(Fe0.9B0.1)1.93 (0<x<0.7) alloys and found that Tb0.2Dy0.4Pr0.4(Fe0.9B0.1)1.93 alloy
with the single Laves phase has a large magnetostriction (λ111=1,200 ppm) and a low anisotropy. This alloy may be a good candidate for magnetostriction applications
BULETIN - APLINDO No.48/2016
23
(a) (b)
Fig. 8: Magnetostriction coefficient vs. magnetic field H (a) and Curie temperature vs. x (b) of alloy
PrxTb0.2Dy0.8-xFe1.85C0.05 (x=0.1-0.4) [15]
3.2 Substitute elements for Fe
3.2.1 Al /Mn
Under low magnetic field, the addition of a small amount of Al can lower the magnetocrystalline anisotropy of the
material, but the magnetostrictive coefficient can be decreased with an increase in Al content. Meanwhile, Curie
temperature will be reduced. In addition, Al is regarded as an ideal substituent for Fe to increase the resistivity
and ductility [19]
. Manganese is an effective substitution element to improve the magnetostrictive property of the
Tb-Dy-Fe alloys. It is noted that the magnetostriction of Mn-containing compounds is larger than that of Mn-free
compounds especially in the lower temperature region. And the addition of Mn can lower the anisotropy energy,
and therefore, a low bias field for saturation magnetostriction is expected. This low bias magnetic field is very
useful since it is sometimes decisive in the practical application [19]
.
3.2.2 Co
The addition of a small amount of Co can stabilize the Laves phase [16]
, but can reduce the magnetostriction
coefficient of materials at the same time [18]
. Replacing Fe by a small amount of Co can increase the alloy’s Curie
temperature TC, but TC will be deceased with the further increase of Co. Z. J. Guo et al. [18]
studied
themagnetostrictive properties of (Tb0.7Dy0.3)Pr0.3(Fe1-xCox)1.85, and the results are shown in the Fig. 9 and Fig.
10, respectively. With increasing Co content, the saturation magnetostriction coefficient decreases, but the Curie
temperature obtains maximum value at x=0.3. As the Co content continues to increase, the Curie temperature
tends to decline.
Z. B. Pan et al. [53]
found that the Co element plays an opposite role in the resultant anisotropy as compared with
Tb. The smallest anisotropy is obtained for the Tb0.3Dy0.6Nd0.1(Fe 0.8Co 0.2)1.93 compound, which has good
magneto-elastic properties, such as the large saturation magnetostrictionλS(~930 ppm) and the high low-field
magnetostrictionλa(~670 ppm/3 kOe).
3.2.3 Si
Eddy current is formed easily in the process of Tb-Dy-Fe alloy in practical applications, which reduces the
efficiency of the transducers. Studies have shown that the eddy current coefficient is inversely proportional to the
electrical resistivity for magnetic material [20]
. Thus, increasing electrical resistivity is a good means
BULETIN - APLINDO No.48/2016
24
Fig. 9: Magnetic field dependence of room temperature magnetostriction λ of annealed
0.025 and 0.1) in 3.5wt.% NaCl aqueous solution. SEM surface morphology
after corrosion test of Tb0.3Dy0.7 Fe1.95 alloy (a), and Tb0.3Dy0.7(Fe0.975Si0.025)1.95 alloy (b)
[54]
BULETIN - APLINDO No.48/2016
26
polarization curves of Tb0.3Dy 0.7(Fe1-xSi x)1.95 (x=0, 0.025 and 0.1) in 3.5wt.% NaCl aqueous solution, and the
SEM surface morphology after corrosion test of Tb0.3Dy0.7Fe1.95 alloy (a), and Tb0.3Dy0.7(Fe0.975Si0.025)1.95 alloy (b).
The surface morphology after corrosion test indicates that the corrosion resistance of x=0.025 is better than that
of the alloy without Si.
3.2.4 Zr
Li Xiaocheng et al. [21]
replaced partial Fe of Tb0.3Dy0.7Fe1.95 alloy by Zr. The addition of different amounts of Zr
(x=0, 0.03, 0.06 and 0.09) has varying effects on alloy magnetostrictive properties. The addition of a small
amount of Zr can effectively restrain the formation of harmful RFe3 phase, which is good for the improvement of
magnetostrictive properties. However, the precipitation of Zr rare earth rich phase is harmful to the
magnetostriction enhancement when x=0.09, which has been shown in Fig. 14.
Fig. 14: Magnetostriction and magnetic field strength curves of alloy Tb0.3Dy0.7Fe1.95-
xZrx (x=0.03, 0.06, 0.09) [21]
Fig. 15: Magnetostriction of Tb0.3Dy0.7(CezFe1-z)1.95 as a function of applied field and
temperature as z=0.75 [22]
3.2.5 Ce
Colm Mac Mahon et al [22]
investigated the magnetization and magnetoelastic properties of melt-spun ribbons of
Tb0.3Dy0.7(CezFe1-z)1.95 (0.025≤z≤0.2). The ribbons exhibit a nanocrystalline structure which becomes more
amorphous with increasingCe content. Room temperature coercivities remain to be 80 kA·m-1
, but low
temperature coercivities increase with the Ce percentage. Saturation magnetostriction varies considerably with
the addition of Ce, reaching a maximum of 850 ppm at 230 K, for z= 0.075 composition as shown in Fig. 15.
4 Heat treatment
The properties of Tb-Dy-Fe alloys are closely related to the material microstructure. After directional solidification,
the Tb-Dy-Fe alloys are usually composed of RFe2 phase and Re-rich phase [38]
. The existence of the Re-rich
phase can improve the toughness of the alloys [55]
. Heat treatment can be used to optimize the morphology of the
BULETIN - APLINDO No.48/2016
27
Re-earth phase, reduce defects, and lower inner stress of the alloys, so that the brittleness of material is
improved [56]
. According to the difference of heat treatment time and procedure, the heat treatment can be divided
into one-step treatment and two-step treatment.
Hu Yong et al. [36]
prepared <110> oriented Tb0.3Dy0.7Fe2 alloy by the method of zone-melting directional
solidification. Results show that the directional solidification Tb-Dy-Fe alloys annealed at 1,203 K for 2 h can
achieve optimal performance with saturation magnetostriction of 1,226 ppm and compressive
Fig. 16: Cleaning tool: (a) Photograph; (b) cleaning station with two devices[66]
strength of 256 MPa. In addition, slow cooling rate can promote high magnetostrictive and mechanical
properties. Chengbao Jiang et al [57]
have successfully prepared <110> oriented rods of TbDyFemagnetostrictive
alloys by zone melting unidirectional solidification. The homogenization annealing for 4 h and 48 h at 1,273 K
have been conducted in a quartz cylinder under Ar atmosphere after pumping to 2×10−3
Pa. A satisfactory
magnetostrictive property of 1,970×10−6
was obtained under 15 MPa pre-stress after heat treatment for 4 h, but
there was not further improvement for 48 h annealing.
Wei Wu et al [55]
have also prepared <110> oriented rods of TbDyFe giant magnetostrictive alloy using zone
melting directional solidification method. Two-step heat treatments were performed at 1,353 K for 2 h, followed
by heating at 673, 773, 873, and 973 K for 4 h in Ar atmosphere and air cooling, respectively. Results showed
that the alloy can get magnetostriction of 1,324 ppm and compressive strength of 585.16 MPa in a magnetic field
of 80 kA·m-1
under 5 MPa pre-stress.
5 Applications
The rare earth giant magnetostrictive material (GMM) is an excellent new functional material. Comparing with
pure nickel and piezoelectric ceramic, Tb-Dy-Fe alloys possess large coupling coefficient and high Curie
temperature as well as higher magnetostriction coefficient [39, 40, 58]
, and have attracted much attention for
applications in high power energy conversion devices [59-61]
. For example, Tb-Dy-Fe alloys can be widely used in
the design of a large-scale ultrasonic cleaning device for boat cleaning [62-67]
, device for high power ultrasonic
spot welding (USW),[68-74]
, and device for therapeutic ultrasound (higherpower ultrasound at lower frequencies) [75,76]
. Moreover, Tb-Dy-Fe alloys also have a potential future in oil exploitation and pipeline transportation [77]
,
and the recycling of waste energy , such as the emulsification and desulfurization of waste tires [87,88]
. Figure
16 exhibits a large-scale ultrasonic cleaning system, and the schematic picture of a multi-transducer device for
boat cleaning (20 kHz). Figure 17 shows the application and the component of the ultrasonic transducer in high-
power ultrasonic oil production. Figure 18 shows the state of the pipeline before installation and six months after
installation of the Tb-Dy-Fe ultrasonic transducer.
BULETIN - APLINDO No.48/2016
28
Fig. 17: Composition of CSYY60H10 high-power
ultrasonic oil production [77]
Fig. 18: State of pipeline (a) before installation and (b) six months after installing Tb-Dy-Fe ultrasonic transducer
6 Conclusion
Giant magnetostrictive material (GMM) is a strategic functional material in the 21st century. Recently, this kind of
material showed a very broad application prospects in military and civilian dual-use high-tech areas. It has
replaced the traditional magnetostrictive materials and has been widely used in advanced technologies, such as
magnetomechanical transducers, actuators and adaptive vibration control systems. As an excellent GMM, Tb-Dy-
Fe alloy possesses large magnetostriction strain, high energy conversion efficiency, and rapid response rate
which have attracted much attention for applications in high power energy conversion devices. However, the cost
of the rare earth element Tb and Dy is too high to be widely applied for the materials. Literatures show that it is
feasible to enhance magnetostrictive properties of the alloy by adding some alloying elements. Nowadays, there
are two different ways to substitute for alloy elements. One is to partially replace Tb or Dy by cheaper rare earth
elements, such as Pr, Nd, Sm and Ho; the other one is using non-rare earth elements, such as Co, Al, Mn, Si,
Ce, B, Be and C, to substitute Fe to form single MgCu2-type Laves phase and a certain amount of Re-rich phase,
which can reduce the brittleness and improve the corrosion resistance of the alloy.
As mentioned above, the properties of the Tb-Dy-Fe alloys play an important role in applications. Therefore, it is
critical to develop new RFe2 compound-based giant-magnetostrictive alloys with excellent properties and lower
cost.
Reference
[1] Clark A E. Magnetic and magnetoelastic properties of highly magnetostrictive rare earth-iron laves phase
compounds. American Institute of Physics Conference Series, 1974, 18: 1015-1029.
[2] Clark AE , Belson HS . Giant Room - Temperature Magnetostrictions in TbFe2 and DyFe2. Physical Review
B, 1972, 5(9): 3642-3644.
[3] Tian Shi. Physical Properties of Materials. Beihang University Press, Beijing, China, 2004: 301. (in Chinese)
[4] Harsh D C and Manfred W. Non-Joulianmagnetostriction. Nature, 2015, 521: 341-343.
[5] Liu XY, Liu JJ, Pan ZB, et al. Optimization on magnetic anisotropy and magnetostriction in TbxHo0.8-
xPr0.2(Fe0.8Co0.2)1.93 compounds. Journal of Magnetism and Magnetic Materials, 2015, 391: 60-64.
[6] Abbundi R, Clark A E. Anomalous thermal expansion and magnetostriction of single crystal Tb0.27Dy0.73Fe2.
IEEE Transactions on Magnetics, 1977, 1(5): 1519-1520.
[7] Jiles D C. The development of highly magnetostrictive rare earth-iron alloys. Journal of Physics D: Applied
Physics, 1994, 27(1): 1-11.
[8] Clark A E. Magnetostrictive rare earth-Fe2 compounds. Handbook of Ferromagnetic Materials, 1980, 1: 531-
589.
[9] Verhoeven J D, Ostenson J E, Gibson E D. The effect of composition and magnetic heat treatment on the
magnetostriction of TbxDy1-xFey twinned single crystals. Journal of Applied Physics, 1989, 66(2): 772-779.
BULETIN - APLINDO No.48/2016
29
[10] Clark A E, Teter J P, McMasters O D. Magnetostriction “jumps” in twinned Tb0.3Dy0.7Fe1.9. Journal of Applied
Physics, 1988, 63(8): 3910-3912.
[11] Clark A E. Magnetostriction in twinned [112] crystal of Tb0.27Dy0.73Fe2. IEEE Transaction on Magnetics, 1986,
22(5): 973-975.
[12] Jiles D C, Thoelke J B. Modelling of the combined effects of stress and anisotropy on the magnetostriction of
Tb0.3Dy0.7Fe2. IEEE Transactions on Magnetics, 1991, 27(6): 5352-5354.
[13] Clark A E, Belson H S, Strakna R E. Elastic properties of rare-earth-iron compounds. Journal of Applied
Physics, 1973, 44(6): 2913-2914.
[14] Wang Bo-wen, Yan Rong-ge. Rare-earth Giant Magnetostrictive Materials, Application and Devices. Journal
of Hebei University of Technology, 2004, 33(2): 16-22.
[15] RenZhi, Li Song-tao, Liu He-yan, et al. Structure and magnetostriction of PrxTb0.2Dy0.8-xFe1.85C0.05 alloys. J Magn Mater Devices, 2013, 44(5): 6-7.
[16] Liu JJ, Pan ZB, Liu XY, et al. Large magnetostriction and direct experimental evidence for anisotropy compensation in Tb0.4-xNdxDy0.6(Fe0.8Co0.2)1.93 Laves compounds. Materials Letters, 2014(137): 274-276.
[17] Wun-Fogle M, Restorff JB, Clark AE, et al. Magnetization and magnetostriction of dendritic [112]
TbxDyyHozFe1.95 (x+y+z=1) rods under compressive stress. Journal of Applied Physics, 1998, 83(11): 7279-
7281.
[18] Guo Z J, Busbridge S C, Wang B W, et al. Structure and Magnetic and Magnetostrictive Properties of
(Tb0.7Dy0.3) 0.7Pr0.3 (Fe1-xCox)1.85(0≤x≤0.6). IEEE Transactions on Magnetics, 2001, 37(4): 3025-3027.
[19] Du J, Wang J H, Tang C C, et al. Magnetostriction in twin-free single crystals TbyDy1-yFe2 with the addition of
aluminum or manganese. Applied Physics Letters, 1998, 72(4): 489-491.
[20] LihongXu, Chengbao Jiang, HuibinXua. Magnetostriction and electrical resistivity of Si doped Tb0.3Dy0.7Fe1.95
[41] S r i s u k h u m b o w o r n c h a i N , G u r u s w a m y S . L a r g e magnetostriction in directionally
solidified FeGa and FeGaAl alloys. Journal of Applied Physics, 2001, 90(11): 5680-5688
[42] Yin H Y, Liu J J, Pan Z B, et al. Magnetostriction of TbxDy0.9- xNd0.1 (Fe0.8Co0.2)1.93 compounds and their
composites (0.20≤x≤0.60). Journal of Alloys and Compounds, 2014, 582: 583-587.
[43] Restorff J B, Wun-Fogle M, Clark A E. Temperature and stress dependences of the magnetostriction in
ternary and quaternary Terfenol alloys. Journal of Applied Physics, 2000, 87(9): 5786-5788.
[44] Wang B, Lv Y, Li G, et al. The magnetostriction and its ratio to hysteresis for Tb-Dy-Ho-Fe alloys. Journal of
Applied Physics, 2014, 115(17): 902-904.
[45] Busbridge S C, Piercy A R. Mannetomechanical properties and anisotropy compensation in quaternary rare
earth-iron materials of the type TbxDyyHozFe2. IEEE Transactions on Magnetics, 1995, 31(6): 4044-4046. [46] Guo Z J, Busbridge S C, Zhang Z D, et al. Microstructure, magnetic properties, and spontaneous
magnetostriction of Tb0.2Pr0.8(Fe0.4Co0.6)x. IEEE Transactions on Magnetics, 2000, 36(5): 3217-3218.
[47] Wang B W. Microstructure and magnetostriction of (Dy0.7Tb0.3)1-xPrxFe1.85 and (Dy0.7Tb0.3)0.7Pr0.3Fey alloys.
Applied Physics Letters, 1996, 69(22): 3429-3431.
[48] Tang Y M, Chen L Y, Zhang L, et al. Temperature dependence of the magnetostriction in polycrystalline
PrFe1.9 and TbFe2 alloys: Experiment and theory. Journal of Applied Physics, 2014, 115(17): 173902.
[49] Yin Hongyun, Liu Jinjun. Research Progress of MgCu2-Type Giant Magnetostrictive Materials with Pr. Rare
Metal Materials and Engineering, 2014, 43(5): 1275-1280.
[50] Ren W J, Zhang Z D, Zhao X G. Magnetostriction and anisotropy compensation in TbxDy1-xPr0.3(Fe0.9B0.1)1.93
[53] Pan Z B, Liu J J, Liu X Y, et al. Structural, magnetic and magnetoelastic properties of Laves – phase Tb 0 .
3Dy0 . 6Nd0 . 1 (Fe 1 - x Co x )1 . 93 compounds (0≤x≤0 . 40) . Intermetallics, 2015, B64: 1-5.
[54] LihongXu, Chengbao Jiang, Chungen Zhou, et al .Magnetostriction and corrosion resistance of Tb0.3Dy0.7(Fe1-xSix)1.95 alloys. Journal of Alloys and Compounds, 2008, 455(1-2):203-206.
[55] Wei Wu, Maocai Zhang, XuexuGao, et al. Effect of two-steps heat treatment on the mechanical properties and magnetostriction of <110> oriented TbDyFe giant magnetostrictive material. Journal of Alloys and Compounds, 2006, 416: 256-260.
[56] Chengbao Jiang, Yan Zhao, LihongXu, et al. Orientation, morphology and magnetostriction of a heat-treated <110> oriented TbDyFe alloy. Journal of Alloys and Compounds, 2004, 373: 167-170.
[57] Jiang C, Zhao Y, Xu L, et al. Orientation, morphology and magnetostriction of a heat-treated<110> oriented TbDyFe alloy. Journal of Alloys and Compounds, 2004, 373(1): 167-170.
[58] Wang Bo-wen, Yan Rong-ge. Rare-earth Giant Magnetostrictive Materials, Application and Devices. Journal of Hebei University of Technology, 2004, 33(2): 16-22.
[59] Jia Z Y, Liu H F, Wang F J, et al. Research on a novel force sensor based on giant magnetostrictive material and its model.
[60] Olabi A G, GrunwaldA . Design and application of magnetostrictive materials. Materials & Design, 2008,
29(2): 469-483.
[61] Joseph M K, Yutang D, Xian Z, et al. Femtosecond Laser Ablated FBG Multitrenches for Magnetic Field
[62] Kubo E, Haibara T, Mori Y, et al. Ultrasonic cleaning method and ultrasonic cleaning apparatus: U.S. Patent
Application 13/892, 327, 2013-5-13.
[63] Niemczewski B. Observations of water cavitation intensity under practical ultrasonic cleaning conditions.
UltrasonicsSonochemistry, 2007, 14(1): 13-18.
[64] Kwan J J, Graham S, Myers R, et al. Ultrasound-induced inertial cavitation from gas-stabilizing nanoparticles.
Physical Review E, 2015, 92(2): 023019.
[65] Eskin G I, Eskin D G. Ultrasonic treatment of light alloy melts. CRC Press, 2014: 32-44.
BULETIN - APLINDO No.48/2016
31
[66] Mazue G, Viennet R, Hihn J Y, et al. Large-scale ultrasonic cleaning system: Design of a multi-transducer
device for boat cleaning (20 kHz). UltrasonicsSonochemistry, 2011, 18(4): 895-900.
[67] Zhimei M. Research Progress in Ultrasonic Scale Inhibition and Elimination. Sino-Global Energy, 2008,
13(4): 92-96. (In Chinese)
[68] Bhosale S B, Pawade R S, Brahmankar P K. Effect of process parameters on MRR, TWR and surface
topography in ultrasonic machining of alumina–zirconia ceramic composite. Ceramics International, 2014,
40(8): 12831-12836.
[69] Liu D F, Cong W L, Pei Z J, et al. A cutting force model for rotary ultrasonic machining of brittle materials.
International Journal of Machine Tools and Manufacture, 2012, 52(1): 77-84.
[70] Panteli A, Robson J D, Brough I, et al. The effect of high strain rate deformation on intermetallic reaction
during ultrasonic welding aluminium to magnesium. Materials Science and Engineering: A, 2012, 556: 31-42.
[71] Panteli A, Chen Y C, Strong D, et al. Optimization of aluminium-to-magnesium ultrasonic spot welding. JOM,
2012, 64(3): 414-420.
[72] Watanabe T, Sakuyama H, Yanagisawa A. Ultrasonic welding between mild steel sheet and Al-Mg alloy
sheet. Journal of Materials Processing Technology, 2009, 209(15): 5475-5480.
[73] Matsuoka S, Imai H. Direct welding of different metals used ultrasonic vibration. Journal of Materials
ProcessingTechnology, 2009, 209(2): 954-960
[74] Matsuoka S. Ultrasonic welding of ceramics/metals using inserts. Journal of Materials Processing
Technology, 1998, 75(1): 259-265.
[75] Mason T J. Therapeutic ultrasound an overview. UltrasonicsSonochemistry, 2011, 18(4): 847-852.
[76] Inoue K, Nakane Y, Michiura T, et al. Ultrasonic scalpel for gastric cancer surgery: a prospective randomized
study. Journal of Gastrointestinal Surgery, 2012, 16(10): 1840-1846.
[77] Wang Z, Xu Y, Suman B. Research status and development trend of ultrasonic oil production technique in
China. UltrasonicsSonochemistry, 2015, 26: 1-8.
[78] Chen T C, Shen Y H, Lee W J, et al. An economic analysis of the continuous ultrasound-assisted oxidative
desulfurization process applied to oil recovered from waste tires. Journal of Cleaner Production, 2013, 39:
129-136.
[79] Adhikari B, De D, Maiti S. Reclamation and recycling of waste rubber. Progress in Polymer Science, 2000,
25(7): 909-948.
[80] Wan M W, Yen T F. Enhance efficiency of tetraoctylammonium fluoride applied to ultrasound-assisted
oxidative desulfurization(UAOD) process.Applied Catalysis A: General, 2007, 319: 237-245.
[81] Quek A, Balasubramanian R. Liquefaction of waste tires by pyrolysis for oil and chemicals-a review. Journal
of Analytical and Applied Pyrolysis, 2013, 101: 1-16.
[82] Holst O, Stenberg B, Christiansson M. Biotechnological possibilities for waste tyre-rubber treatment.
Biodegradation, 1998, 9(3-4): 301-310.
[83] Al-Lal A M, Bolonio D, Llamas A, et al. Desulfurization of pyrolysis fuels obtained from waste: Lube oils, tires
and plastics. Fuel, 2015, 150: 208-216.
[84] Liu L, Wen J, Yang Y, et al. Ultrasound field distribution and ultrasonic oxidation desulfurization efficiency.
UltrasonicsSonochemistry, 2013, 20(2): 696-702.
[85] Chen T C, Shen Y H, Lee W J, et al. The study of ultrasound-assisted oxidative desulfurization process
applied to the utilization of pyrolysis oil from waste tires. Journal of Cleaner Production, 2010, 18(18): 1850-
1858.
[86] Wan M W, Yen T F. Portable continuous ultrasound-assisted oxidative desulfurization unit for marine gas oil.
Energy & Fuels, 2008, 22(2): 1130-1135.
[87] Yao Xiu-qing, Zhang Jie, Li Fei-fei, et al. Recent Process of Desulfurization Technology of the Clean Fuel.
Journal of Liaoning University of Petrol EUM and Chemical Technology, 2004, 24(1): 39-42.1
[88] Dong Chengchun . A brief introduction of ultrasonic desulfurization. Rubber & Plastics Resources Utilization,
2012(3): 27-29. (In Chinese)
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BULETIN - APLINDO No.48/2016
32
Effects of Si alloying and T6 treatment
on mechanical roperties and wear
resistance of ZA27 alloys
Rui Zhang, Guang-lei Liu, *Nai-chao Si, Yu-yang Peng, Hao Wan, and Ting Liu
School of Materials Science and Engineering, Jiangsu University, Zhenjiang 212013, China
Abstract: To improve the mechanical properties and wear resistance of ZA27 alloy, Si was introduced to
thealloy, and the effect of Si alloying and T6 heat treatment on the microstructure, mechanical properties and
wear resistance was investigated. The results show that with 0.55% Si, the microstructure of the alloy can be
refined effectively, which leads to the increase of hardness. But the tensile strength and elongation decrease
because Si undermines the integrity of the matrix. On the other hand, the dendrites are transformed into a
desired α+η+(α+η) mixture with T6 heat treatment, which introduces a remarkable increase to the elongation and
hardness of the alloy. The wear resistance of the ZA27 alloy with Si alloying is significantly better than that of the
ZA27 alloy without Si. With the increase of Si addition, the wear resistance of the alloy firstly increases and then
decreases. In the alloy without Si alloying, severe plastic deformation and large delamination were observed on
the worn surface of the alloy. However, with the increase of Si, the main wear mechanism transformed to
abrasive wear gradually. In addition, the T6 treatment can further improve the wear resistance of the alloy with Si
alloying.
Key words: ZA27 alloy; Si alloying; mechanical properties; wear resistance
CLC numbers: TG146.21 Document code: A Article ID: 1672-6421(2016)02-093-08
As-cast zinc-aluminum alloy has been developedfrom late 1930s, which attracted attention of researchers for
decades as a promising material [1-3]
. The alloy has been widely applied to various fields. One of the most
important applications of ZA alloy is as wear parts under low-speed heavy-duty conditions, as a substitute for tin-
bronze due to its better wear resistance, lower cost and longer service life [4-6]
. ZA alloys show advantages in
mechanical properties as compared with traditional non-ferrous alloy. The study by Chen T J, et al [7]
revealed
that ZA alloys have lower friction coefficient and higher bearing capacity than traditional wear resistant materials.
The friction coefficient of the ZA27 alloy is even lower than copper alloys through complex modification with RE,
Ti, B and Zr [8]
. However, composition segregation, poor dimensional and property stability are the main
disadvantages, limiting the application of ZA alloy in modern industry. To extend its application area, many
optimized processes are used to improve and balance the properties.
In recent years, many new effective alloying elements (Cu, Mn, Ti, Re, Si) [9-13]
and alloying methods were
discovered, which can improve the mechanical properties and wear resistance of ZA alloys. For example, with
0.4% Ni addition, the microstructures of ZA27 alloys were refined effectively and the wear resistance under high -
speed heavy - duty conditions was significantly improved [14]
. Li Zi-quan and Zhou Heng-zhi [15]
investigated the
microstructure characteristics of aged SiCp/ZA27 composite, and their study results demonstrated that SiC
particulates strongly accelerate neighboring β phase decomposition in the aging process. Stabilizing and
solution-aging treatments were typically used in the heat treatment of ZA alloys for refinement, stability and
homogenization of the microstructures [16]
. Almost all the previous studies were involved in single optimizing
process. Very few literatures could be found focusing on composite process for ZA27 alloys. In this paper, Si
alloying and T6 heat treatment were used for improving the mechanical properties and wear resistance of the
ZA27 alloy. The results will provide a basis for the complex treatment of ZA alloys.
* Nai-chao Si
Male, born in 1956, Professor, Ph.D supervisor. His research interests mainly focuse on seismic and vibration damping performance in engineering structure of Cu based shape memory alloys, application of high strength thin walled gray cast irons and austempered ductile irons in automobile engine; and performance optimization of nonferrous alloys.
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1 Experimental procedure
1.1 Alloy preparation
The nominal compositions of the ZA27 alloy (in wt.%) are shown in Table 1. Silicon addition in wt.% was 0, 0.3,
0.55 and 0.8, respectively. The alloy was made from commercial purity aluminum (99.80%), zinc (99.99%),
magnesium (99.5%), Al-50wt.%Cu master alloy and Al-7wt.%Si master alloy. The aluminum was melted at 700
°C at first, and then the Al-50wt.%Cu and Al-7wt.%Si master alloys were added into the melt. After the master
alloys were melted, the zinc was added into the melt. Mechanical mixing for 15 min through a stainless steel
stirrer coated with aluminite was applied to ensure homogeneous distribution of the elements in the melt. Then
magnesium was pressed into the bottom of the melt to reduce the amount of burning loss. After 5 min, the
C2H2Cl6 agent was bubbled into the melt for degassing. Then the melt was refined with 0.2% dewatered ZnCl2 for
10 min. The overheated melt (600 °C) was cast into a preheated columnar steel mold (200 °C) to obtain alloy
samples (Φ35 mm × 270 mm). Wear samples (20 mm × 10 mm × 8 mm) were fabricated using a Wire-Electronic
Discharging Machine and tensile samples through machining. One group of the specimens were subjected to
heat treatment of solution at 365 °C for 6 h, then quenched in water and artificially aged at 160 °C for 4 h (T6).
Table 1: Nominal chemical compositions of ZA27 alloy (wt.%)
1.2 Measurement of mechanical properties and microstructural characterization
Tensile tests were carried out at room temperature on a 600 kN hydraulic universal testing machine (WE-600) at
a 3 mm·min-1
tensile rate. Dimensions of the tensile bar are shown in Fig. 1. Three sets of measured data were
used to calculate averages. Bulk hardness of all samples was measured using a Brinell hardness tester with a 5
mm diameter steel ball indenter and under a load of 2.452 kN. The measured impression diameter was used in
equation 1 for calculation.
Where F is the load, D is the diameter of steel ball, and d is the indentation
Microstructures of corroded surfaces of the samples were observed under a NIKONPIPHOT300 optical
microscope. The corrosives applied consisted of diluted hydrochloric acid (1 vol.%), dilute nitric acid (1 vol.% ),
diluted hydrofluoric acid (2 vol.%) and distilled water (96 vol.%)[17]
.
Al Cu Mg Zn
26-28 2.0-2.5 0.030-0.04 Balance
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1.3 Sliding wear tests
Wear tests were carried out on a block-on-disc friction and wear tester (M-2000). Figure 2 shows the operating principle of the sliding wear process. The wear test cycle lasted 3 h under the load of 600 N with a rotational speed of 200 r·min
-1, and the friction counterpart was made of GCr15. Lubrication was provided by dropping
lubricating oil SAE 30 onto the friction surface of the rotating disk at a rate of 15 to 20 drops per min. Wear mass loss was calculated by the difference in sample weight measured before and after the wear test. Coefficients of friction were recorded per min from 30 min after the test start to the end. The coefficient of friction was calculated by equation 2.
where T is time, r the radius of circle, b the width of worn surface, p the load, and θ is equal
2 Results
2.1 Mechanical properties
Mechanical properties of the ZA27 alloys with different contents of silicon, and in both as-cast and heat-treated conditions, are shown in Table 2. The increase of Si content caused a slight decrease of the tensile strength and elongation, while their hardness increased with the increase of Si%. A remarkable increase of the elongation and a decrease of the tensile strength were caused by T6 heat treatment. In addition, the hardness of heat treated samples slightly increased compared to that of the as-cast alloy.
2.2 Microstructure
Microstructures of the as-cast alloys are shown in Fig. 3. In the alloy without Si alloying, substantial amounts of large dendritic crystals, developed second dendrite arms and bits of third dendrite arms can be observed (Fig. 3a). However, in the alloy with 0.3% Si, large dendritic crystals decreased and second dendrite arms reduced and shortened, as shown in Fig. 3b. It can be seen in Fig. 3c that, as the content of Si reached 0.55%, almost all the dendrites transformed into equiaxial snowflake grains or blocky crystals. However, coarse dendrites appeared again with 0.8wt.% Si, as shown in Fig. 3d. It reveals that a certain amount of silicon can refine the microstructure of ZA27 alloy. In the process of solidification, the constitutional supercooling formed in the front of solid-liquid interface due to the enrichment of Si, resulted in branches necking and fusing in the process of crystal
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Fig. 3: Microstructure of as-cast alloy with Si content (wt.%): (a) 0, (b) 0.3, (c) 0.55, (d) 0.8
growth. In other words, the growth of α-dendrites was prevented from the enrichment of Si. At the same time, the
growth of separated grains was promoted with temperature-fluctuation which benefited the refinement of grains.
However, excessive Si would precipitate as primary Si and reduce the content of Si in melt, which weakened the
effect of constitutional supercooling and resulted in coarsening of grains.
Figures 4a-4c show magnified microstructures of the ZA-27 alloys containing 0.3wt.%, 0.55wt.% and 0.8wt.% Si,
residual η phase, ε phase and some black phases in the interdendritic regions. (A- α dendrites core, B- α+η
eutectoid structure, C- black phase, D-ε phase). Spectral analysis was carried out on the black phase in Fig. 4c,
and the result is shown in Fig. 4d, identifying that the marked zone was primary Si phase. It can also be seen that
the morphology of primary Si changed from rod-like to blocky with the increase of Si addition.
The microstructure of the ZA27 alloy with 0.55wt.% Si in solution and aging treated condition is shown in Fig. 5.
Granular zinc-rich η phases were uniformly distributed in gray matrix structures instead of dendritic structure.
After solution treatment at 350 °C, the matrix was β phase. When the specimens were quenched in water (70
°C), part of β phases transformed into (α+η) phase through eutectoid reaction, and a large proportion of β phases
retained. When aged at 160 °C, supersaturated η phase precipitated from the residual β phases. The matrix
transformed into a α+η mixture. The magnified microstructure in Fig. 5b shows two different mixtures, the lamellar
structure produced by eutectoid reaction and the small spherical mixture produced by aging. Through aging heat
treatment, η phase was formed by zinc enrichment in local area, and other zinc elements were dispersively
distributed in the matrix structure. With the combined effects, the microstructure of fine η phase and α+η mixture
was produced by the heat treatment.
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2.3 Wear properties
The effect of T6 heat treatment on wear loss of ZA27 alloys with different silicon additions is shown in Fig. 6. It can be noticed that wear loss reached the maximum when silicon content was zero and wear loss of the as-cast alloy with 0.55wt.% Si alloying reached the minimum. Thus, Si alloying appears to do the best optimization for wear resistance of ZA27 alloy when Si was 0.55wt.%. Wear losses of T6 heat-treated alloys decreased drastically. Similarly, wear loss of the heat-treated alloy reached the minimum with 0.55wt.% Si.
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Figure 7 shows worn surfaces of the as-cast ZA27 alloy. Figure 7a is the worn surface of the alloy without Si alloying. Distorted polishing scratches, sags and crests can be found on the surface which may be caused by friction heat. In Fig. 7b (0.3wt.% Si), large-scale of delamination was observed on the worn surface, which could be attributed to adhesive wear mechanism. Continuous scratches appeared on the surface when Si% increased up to 0.55wt.%, which was caused by abrasive wear mechanism (Fig. 7c). Although adhesive wear still existed, the extent was greatly reduced, and the worn surface became relatively smooth. Abrasive wear became the main wear mechanism. When Si addition increased to 0.8wt.%, adhesive
wear became aggravated (Fig. 7d), because blocky or rod-shaped primary Si phases had undermined the integrity of the matrix alloy. In this case, Si phase can be separated by friction force along the direction perpendicular to the force. Stress concentrations arising in these small gaps led to the formation of cracks and delamination.
Worn surfaces of the T6 heat-treated alloys are shown in Fig. 8. The worn surface of the alloy without Si alloying is displayed in Fig. 8a. Slight adhesive wear occurred on the worn surface. Delamination still existed on the worn surface of the alloy without Si alloying, but the thickness and size decreased significantly. Abrasive wear mechanism became the main wear mechanism of the alloy with 0.3% Si (Fig. 8b). But polishing scratches were still thick and broad. As the Si% increased to 0.55% (Fig. 8c), abrasive wear became the predominant wear mechanism and the worn surface tended to be smooth and clean with fine polishing scratches. When the Si addition reached 0.8wt.% (Fig. 8d), abundant blocky Si phase formed, which could easily split away from the matrix. The blocky Si phase could serve as wear debris to cut the matrix, leading to relatively thick scratches again.
Fig. 7: Wear appearances of as-cast alloys with Si addition (wt.%): (a) 0, (b) 0.3, (c) 0.55, (d) 0.8
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Fig. 8: Worn surfaces of heat-treated alloys with Si addition (wt.%): (a) 0, (b) 0.3, (c) 0.55, (d) 0.8
3 Discussion
The effects of Si alloying on mechanical properties of the ZA27 alloys are shown in Table 2. The results indicate
that the addition of Si reduced the tensile strength, while increased the hardness. Si alloying refined the
microstructure of the alloy (Fig. 3), but at the same time, primary Si phases undermined the integrity of the matrix
alloy (Fig. 4a, 4b, 4c). The morphologies of primary Si phases are usually rod-shaped and blocky. The small-
angle gap in the matrix alloy formed by sharp ends of Si phase easily produces stress concentration, leading to
the decrease of the tensile strength and elongation. But the hardness should not be impacted by this effect. On
the contrary, the hardness of the alloy increased due to the high hardness silicon crystal and the refined
microstructure.
Through the T6 heat treatment, the tensile strength decreased, but the elongation increased significantly (Table
2). This result is in line with Babic Miroslav’s research [18]
. A new α + η + (α+η) microstructure was reformed by
the heat treatment.
Deformation on the mixture was more uniform, which increased the elongation. The soft η phase was dispersively
distributed on the matrix, decreasing the difficulty of deformation, which led to the decrease of deformation force.
The friction coefficient of the as-cast ZA27 alloys during the sliding wear test is shown in Fig. 9. Since the initial
30 min was the running-in period, records of friction coefficients started from 30-min mark. From the diagram, it
can be clearly noticed that high friction coefficient and drastic fluctuation occurred in the curve of the ZA27 alloy
without Si alloying (Curve-A)
because of adhesive wear. With the increase in Si content, the average friction coefficient decreased significantly
(Curve-B, Curve-C). But when Si content reached 0.8wt.%%, friction coefficient of the alloy (Curve-D) increased
to 0.04 on average, and fluctuation of the friction coefficient increased. This is in accordance with Fig. 7.
In Fig. 6, the T6 heat-treated ZA27 alloy with 0.55wt.% Si shows the best wear resistance. Silicon particles and
ε(CuZn4) phases acted as supporting load and limited the direct contact between Zn-Al matrix and the steel
slider. At the microcosmic level, soft matrix phases were first worn off, and hard spots highlighted on the wear
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contact surface and acted as the supporting load, avoiding abrasions of the soft matrix. At the same time, the
worn parts played a good storage function in the
Fig. 9: Friction coefficient of as-cast alloys
condition of oil lubrication. The ε(CuZn4) phase showed the feature of high hardness, but due to its small
proportion in the matrix, it cannot support the heavy load perfectly, and may even have the opposite effect.
Namely, the ε(CuZn4) phase may be forced to cut into the matrix. The situation happened in the wear process of
the ZA27 alloy without Si alloying. The temperature on surface of the alloy increased rapidly due to the direct
contact between the matrix and the steel slider, resulting in the softening effect of the surface layers. Adhesive
wear mechanism became an important wear mechanism, leading to the drastic fluctuation of friction coefficient.
When excessive Si was added, primary rod-shaped Si phases changed to blocky particles. The blocky Si
particles could easily split away from the matrix, and then became abrasive particles to cut the alloy matrix. Thus
the wear property deteriorated when Si increased to 0.8wt.%.
The characteristics of the friction coefficient of T6 heat-treated
ZA-27 alloys during sliding are illustrated in Fig. 10. The
monolithic friction coefficient and the fluctuation decreased
enormously compared with the as-cast alloy. The friction
coefficient of the heat treated alloy without Si alloying was
about 0.042. It reached the minimum (about 0.015) when Si
content was 0.55%, while with further increase of Si addition,
the friction coefficient increased again. It can be seen that the
reduction in friction coefficient occurred in the rear part of
curves C and D, which may be caused by surface hardening.
T6 heat treatment transformed the microstructure into a fine
mixture. Under the condition of oil lubrication, small spherical
η phases were worn away first, then the pits in reserve were
filled with lubricated oil. This structure improved the wear
resistance of the alloy.
Figure 11 shows the deformation on the edge of the alloy specimen without Si alloying. The lamellar structure
was caused by extrusion force, which was the product of stress fatigue.
The wear debris is shown in Fig. 12. The debris of the alloy without Si alloying is displayed in Fig. 12a,
presenting water ripples on the surface. The wear process can be repeated based on all the above facts
presented in Figs. 7a, 11 and 12a: micro-cracks in the subsurface formed under the shear stress, and
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extended parallelly in a certain depth from the surface, creating a gap between the surface and subsurface. Due to the plastic deformation and instantaneous high temperature caused by local high press, cold welding spot formed between the steel slider and the alloy surface. Along with the slide of the grinding wheel, the local surface of the alloy peeled off. It can be concluded that adhesive wear and fatigue wear are the main wear mechanism of the alloy without Si alloying. The temperature on the surface of the alloy increased rapidly due to the direct contact between the matrix and the steel slider, resulting in the softening effect of the affected layers. The debris of the alloy with
0.55wt.% Si is shown in Fig. 12b. Granulated particles, like globular, cubic or other shapes were the products of
abrasive wear. In addition, the dimension of the debris decreased obviously. At this time, abrasive wear became
the main wear mechanism.
4 Conclusions 1) The microstructure refined by Si alloying is the main reason for the increase of hardness. Meanwhile the
integrity of the matrix undermined by Si alloying causes the decrease of the tensile strength and elongation.
2) A α + η + (α+η) mixture formed through T6 heat treatment causes the decrease of the tensile strength. As
compared to the as-cast alloy, the heat-treated samples obtain remarkable increase of elongation and
hardness.
3) Wear resistance of both the as-cast and T6 treated alloys firstly increases and then decreases with the
increase of Si content. With same Si content, wear resistance of the T6 alloy is better than the as-cast. When
Si addition reaches 0.55wt.%, wear resistance achieves the best for both as-cast and T6 treated alloys.
Fig. 12: Wear debris of alloys (0% Si in as-cast
and 0.55wt.% Si in T6 temper)
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4) For the as-cast alloy, wear mechanism transforms from adhesive wear and fatigue wear into abrasive wear
with the increase of Si content. T6 heat treatment is beneficial to the wear resistance, and abrasive wear is
the main wear mechanism.
References
[1] Sastry S, Krishna M, Uchil J. A study on damping behaviour of aluminite particulate reinforced ZA-27 alloy metal matrix composites. Journal of Alloys and Compounds, 2001, 314(1):268-274.
[2] Li Yuan-dong, Zhang Xin-long, Ma Ying, et al. Effect of mixing rate and temperature on primary Si phase of hypereutectic Al-20Si alloy during controlled diffusion solidification (CDS) process. China Foundry, 2015, 12(3): 173-179.
[3] Geng Hao-ran, Tian Xian-fa, Cui Hong-wei, et al. Antifriction and wear behaviour of ZAS35 zinc alloy. Influence of heat treatment and melting technique. Materials Science and Engineering A, 2001, 361: 109-114.
[4] Chen Fei, Wang Tong - min, Chen Zong - ning, et al . Microstructure, mechanical properties and wear behaviour of Zn-Al-Cu-TiB2 in situ composites. Trans. Nonferrous Met. Soc. China, 2015, 25: 103-111.
[5] Bobic Biljana, Bajat Jelena, Aimovic-Pavlovic Zagorka, et al. Corrosion behaviour of thixoformed and heat-treated ZA27 alloys in NaCl solution. Trans. Nonferrous Met. Soc. China,2013, 23: 931−941.
[6] Chen T J, Hao Y, Sun J. The microstructural and constitutional evolution of cast dendritic ZA27 alloy during partial remelting. Journal of Materials Processing Technology. 2004, 148: 8-14.
[7] Chen T J, Hao Y, Sun J, et al. Effects of processing parameters on tensile properties and hardness of thixoformed ZA27 alloy. Materials Science and Engineering A, 2004, 382: 90-103.
[8] Tan Yinyuan. Effects of compound modifier on microstructure and performance of ZA27 alloy. Journal of Nanjing University of Science and Technology, 2002, 05: 547-551.
[9] Zhu Y H, Man H C, Dorantes-Rosales H J, et al. Ageing characteristics of furnace cooledeutectoid Zn-Al based alloy. Journal of Materials Science, 2003, 38: 2925-2934.
[10] Zuo Yu-bo, Liu Xu-dong, Sun Chao, et al. Grain refinement and macrosegregation behavior of direct chill cast Al-Zn-Mg-Cu alloy under combined electromagnetic fields. China Foundry, 2015, 12(5): 333-338.
[11] Chen Ti-jun, Li Yuan-dong, Hao Yuan. Effects of Mg and RE additions on the semi-solid microstructure of a zinc alloy ZA27.Science and Technology of Advanced Materials, 2003, 4(6):495-502.
[12] Xu Xiao-qing, Li Dr-fu, Guo Sheng-li, et al. Microstructure evolution of Zn-8Cu-0.3Ti alloy during hot deformation. Transactions of Nonferrous Metals Society of China, 2012,22(7): 1606−1612.
[13] Chen Ti-jun, Zhang Da-hua, Wang Wei, et al. Effects of Y content on microstructures and mechanical properties of as-cast Mg-Zn-Nd alloys. China Foundry, 2015, 12(5): 339-348.
[14] Wang Huai-qing, Si Nai-chao, Si Song-hai, et al. Effect of Ni Alloying on Microstructure and Wear of ZA27 Alloy. Tribiology,2013, 33(1): 57-64.
[15] Li Zi - quan, Zhou Heng - zhi, Luo Xin - yi, et al . Aging microstructural characteristics of ZA-27 alloy and SiCp/ZA-27 composite. Trans. Nonferrous Met. Soc. China, 2006, 16: 98-104.
[16] Liu Yang, Li Hong-ying, Jiang Hao-fan, et al. Effects of heat treatment on microstructure and mechanical properties of ZA27 alloy. Trans. Nonferrous Met. Soc. China, 2013, 23: 642-649.
[17] Wu Yong-yong, Si Nai-chao, Liu Guang-lei, et al. Effect of Mn Alloying on Microstructures and Wear Property of ZA43 Alloy.Foundry, 2014(10): 1019-1023.
[18] Babic Miroslav, Aleksandar Vencl, Slobodan Mitrovic, et al.Influence of T4 Heat Treatment on Tribological Behavior of Za27 Alloy Under Lubricated Sliding Condition. Tribol Lett, 2009, 36: 125-134.
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Effects of grain refinement on cast structure and tensile properties of superalloy K4169 at high pouring temperature Zi-qi Jie 1, Jun Zhang 1, *Tai-wen Huang 1, Lin Liu 1, Hai-jun Su 1, Yan-li Shi 2, and Heng-zhi Fu 1
1. State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi’an 710072, China
2. Xi’an Jiaotong University City College, Xi’an 710072, China
Abstract: In order to improve the filling ability of large complex thin wall castings, the pouring temperature should be increased, but this will result in the grain coarsening. To overcome this problem, two kinds of grain refiners of Co-Fe-Nb and Cr-Fe-Nb ternary alloys, which contain high stability compound particles, were prepared. The effects of the refiners on the as-cast structures and tensile properties of the K4169 superalloy with different casting conditions were studied by analyzing specimens 110 mm long and 20 mm in diameter. Results showed that the mixture addition of the two refiners in the melt of K4169 can reduce the columnar grain region and decrease the equiaxed grain size greatly. After refinement, the amount of Laves phase decreases and its morphology changes from island to blocky structure. The carbides in the fine grain samples are fine and dispersive. Meanwhile, the porosity in specimens is decreased due to grain refinement. As a result, the yield strength, ultimate strength and the elongation of the specimens are increased. The grain refinement mechanisms are also discussed.
Key words : superalloy; K4169; grain refinement; tensile properties
CLC numbers: TG143.9 Document code: A Article ID: 1672-6421(2016)02-101-06
With the continuous improvement of the engine thrust-weight ratio, the turbine disk and the intermediate case in the turbine engine become more complex in structure. Meanwhile, the operating temperature of these superalloy castings reaches approximately 700 °C. Under such a high temperature, a uniform and fine-grained microstructure is desirable in order to obtain good low cycle fatigue resistance and high tensile strength
[1, 2].
Therefore, the production of such components challenges metallurgists and requires the development of advanced casting technology. One of the effective methods of improving the filling ability of large-size thin-wall castings is to increase the pouring temperature, but this will lead to a coarse grain size and decrease the low cycle fatigue resistance and tensile strength. The addition of the grain refiners in the melt is an effective way to overcome the problem and to get fine grain size
[1].
Grain refinement of as-cast structure means increasing the heterogeneous nucleation sites during the solidification of castings. Fine grain casting techniques of superalloys mainly include the thermal control method
[3], the
chemical approach [1, 4, 5]
and the dynamic method [6-8]
. Among these, the chemical fine grain process is an efficient method, where the heterogeneous nucleation can be increased by the addition of a specially designed master alloy, which contains suitable solid particles with high stability in the melt. This method needs neither complicated equipment nor complex process. Refractory metal oxides, carbides, nitrides and boron have been used as refiners in some superalloys
[1, 9]. However, this kind of refiner will introduce inclusions in the castings, which may become
crack initiation sites, and deteriorate the mechanical properties [10]
. Especially, the addition of boron will decrease incipient melting temperature of the alloys, which reduces the plastic properties greatly
[11]. Liu et al
[1] developed
two kinds of refiners Co-Fe-Nb and Cr-Mo-Nb used in Ni-Fe based super alloys. These refiners possess effective refinement capability without introducing inclusions. However, Co-Fe-Nb and Cr-Mo-Nb can only be used at temperatures of 1,360-1,420 °C, far below the melting and pouring temperatures for most superalloy castings.
Ni-based superalloy K4169 is widely used in turbine disk and intermediate case components, due to its high-temperature mechanical properties in addition to an excellent corrosion resistance
[12, 13]. While the shape of
these components tends to become more complex and thin-walled, leading to the bad filling. In order to improve the filling ability of complex and thin-walled castings, the pouring temperature should be increased, but this will result in grain size coarsening. To obtain good filling ability and fine microstructure, two inter-metallic compounds Co3FeNb2 and CrFeNb were prepared as refiners of the K4169 alloy. The constituent elements are also the elements presented in the superalloy K4169, ensuring that the grain refiners do not introduce inclusions and pose any harmful influence on the mechanical properties of the alloy. The effect of grain refiner on cast structure and tensile properties of K4169 at the pouring temperatures of 1,470-1,520 °C was investigated.
* Tai-wen Huang Male, born in 1975, Ph. D., Professor. His research interests mainly focus on superalloys. E-mail: [email protected]
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Fig. 1: Grain structures for different treatment conditions and pouring temperatures: (a) without refiner addition, 1,520 °C,
(b) without refiner addition, 1,470 °C; (c) with refiner addition, 1,520 °C; (d) with refiner addition, 1,470 °C
1. Materials and experimental procedure Two kinds of ternary alloy grain refiners with the nominal compositions of Co3FeNb2 and CrFeNb were designed
and the button ingots of the refiners were prepared by melting an appropriate proportion of the constituents in an
arc melting furnace in an argon atmosphere. The raw materials were 99.95% Cr powder, 99.9% Co powder,
99.5% Fe powder and 99.97% Nb block. They were ground into powders with a size of 60-100 μm. The physical
and crystallographic parameters of the refiners are listed in Table 1. The melting point of the refiners was
analyzed by differential thermal analysis (DTA). The mixture of the two refiners was prepared by mixing
physically with the proportion of 1:1 in weight percentage.
Table 1: Physical and crystallographic parameters of experimental refiners
Refiner Crystal
structure Density (g·cm-3
) Melting point
(°C)
Co3FeNb2 Hexagonal 8.8 1,550
CrFeNb Hexagonal 8.2 ﹥1,650
The commercial K4169 alloy with the composition (wt.%): 0.056 C, 0.01 Co, 52.54 Ni, 19.15 Cr, 3.11 Mo, 0.61 Al,
0.94 Ti, 5.03 Nb, 0.0026 B, 0.028 Zr with the balance being Fe was used for the grain refinement experiments.
The equilibrium liquidus and solidus temperatures of the alloy are 1,349 and 1,270 °C, respectively, according to
DTA results.
A vacuum melting furnace was used to cast ingots of K4169 superalloy. The melt was first superheated up to
1,550 °C and held for 2-4 min and then cooled down to the pouring temperature. For the conventional cast
samples, the melt was poured into the preheated mold directly. However, for chemical grain refinement samples,
the refiner was added into the melt at the pouring temperature. The addition amount was 0.3wt.% of the charge.
After that, the melt was stirred for the refiner particles to be dispersed in the melt uniformly. Then the melt was
held for 30-60 s for homogenization of the refiner and subsequently poured into the mould. The ceramic moulds
with inner size of 120 mm in length and 20 mm in diameter and the preheating temperature of 900 °C were used
in all cases.
The as-cast ingots were sectioned along the cross-section and the samples were ground, polished and
subsequently chemically etched with a solution etchant of 15 g CuSO4, 3.5 ml H2SO4 and 50 ml HCl to expose
grain structures. The average equiaxed grain size and fraction of equiaxed grains at transverse cross-section
were determined by a standard quantitative metallographic technique. The grain size was measured by the line
intercept method and estimated with reference to the ASTM standard. The distribution of the alloying elements
was determined using Electron-probe microanalysis (EPMA). The tensile tests were conducted using an Instron
3382 testing machine at room temperature. At least three identical specimens were tested for each case.
2. Results
The structures of transverse sections of different as-cast samples are shown in Fig. 1. For the samples without
grain refinement, only columnar grains were observed. By adding the refiners, an equiaxed grain region was
formed and the average grain sizes were reduced. For the samples with the pouring temperature of 1,520 °C,
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after the addition of mixture refiners of the two ternary inter-metallic compounds, the average grain size was refined from 10.56 to 2.84 mm and the proportion of equiaxed grains at cross-section was increased from 10% to 81%. Similar results were obtained for the samples with the pouring temperature of 1,470 °C, where the average grain size was refined from 8.98 to 1.85 mm and the proportion of equiaxed grains was increased from 15% to 93%. The dendritic morphologies are shown in Fig. 2. It can be seen that the dendritic morphologies with highly developed branches were obtained in the case without grain refiner
addition. However, the average length of the primary dendrite axes decreases with the addition of the grain
refiners. The secondary dendrite arm spacing (SDAS) of both samples (a) and (b) was about 65 μm, and for (c)
and (d) was about 58 μm, indicating that the grain refinement has a negligible effect on the SDAS, but the SDAS
decreases with the decrease of the pouring temperature.
K4169 superalloy has a wide solidification temperature range. Therefore, porosity is likely to form in its castings. Figure 3 shows the morphology and the distribution of porosity in the samples with the grain sizes of 10.56 mm and 1.85 mm. It was shown that there exist some intensively distributed large-sized porosities in coarse grain samples. However, it becomes uniform and much smaller in the chemically refined specimen
Fig. 3: Porosity of samples with different casting conditions: (a) grain size, 10.56 mm; (b) grain size, 1.85 mm
Fig.2 : Dendritic morphologies for different casting conditions: (a) without refiner addition, 1,520 °C; (b) with addition, 1,520 °C; (c) without refiner addition, 1,470 °C; (d) with addition, 1,470 °C
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The typical as-cast microstructure of K4169 consists of the primary gamma phase dendrites, carbides, laves and delta phase
[13]. Micrographs of laves and MC carbides in test bars of different grain sizes were obtained. In the
test bars with the grain size of 10.56 mm and 2.84 mm, block carbides and eutectic laves can be observed. The carbide morphology in the fine grain samples is fine and dispersive. However, the laves phase is mainly contained in the eutectic phase in the coarse grain.
Fig. 4: Microstructure of alloy with different casting conditions: (a) without grain refiner addition, 1,520 °C,
grain size 10.56 mm, (b) with refinement, 1,520 °C, grain size 2.84 mm
Figure 5 shows the correlation of grain size and ultimate tensile strength and yield strength obtained at the room temperature tensile tests. The ultimate tensile strength and yield strength of K4169 superalloy are significantly
improved along with the grain refinement. When the grain size of K4169 superalloy is decreased from 10.56 mm to 2.84 mm at the pouring temperature of 1,520 °C, the tensile and yield strength are increased by 11.76% and 9.8%,
respectively. For the pouring temperature of 1,470 °C, the tensile and yield strength are increased by 19.07% and
29.16%, respectively, corresponding to the grain size decreases from 8.98 mm to 1.85 mm. In addition, the elongation is increased with the addition of grain refiners at different pouring temperatures. At the pouring temperature of 1,520
°C, when the grain size is samples, and the block laves are found in the fine grain samples. The quantity of laves
decreases with the decrease of grain size. Besides, the quantity of carbides remained about the same for the same pouring temperature. The results in Fig. 4 show that the volume fraction of laves phase is about 3.35% when the grain size is 10.56 mm. It was reduced to 1.48 % if the grain is refined to 2.84 mm. refined from 10.56 mm to 2.84 mm, the elongation is increased by 53%. When the pouring temperature is 1,470 °C, the elongation is increased by 38% corresponding to the grain size decreases from 8.98 mm to 1.85 mm.
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3 Discussion
Results of this study show that the refiners can lead to grain refinement and increase the proportion of equiaxed
grains. The main principle is a fine epitaxial fit between low-index planes of the heterogeneous nucleation particle
substrate offered by the grain refiners and the nucleated solid phase. The lower the lattice disregistry, the more
effective the refiner will be in promoting nucleation. According to the calculation model of lattice disregistry (δ)
between refiners and the nucleated phase proposed by Bramfitt [14]
, when the value of δ for some specific crystal
planes is less than 12%, the refiner will have a good refining effect. Wang et al [15]
calculated and simulated the
planes matching models and matching orientations. The results show that (0001) and (0110) planes of refiners
Co3FeNb2 and CrFeNb have a fine crystallographic matching relationship with the (110), (111) planes of γ matrix
of K4169. Therefore, the refiners can act as the nucleation substrate of γ matrix and allow its epitaxial growth.
Presence of a great number of active refiner particles in the melt would cause enormous heterogeneous nuclei of
crystallites, which would impinge on one another and restrict further growth. Hence, the formation of numerous
nuclei and the restriction on their further growth result in the refinement of grains. However, due to the higher
pouring temperature, the refining effect is reduced.
It can be seen from Fig. 1 that adding refiner to the melt makes the equiaxed fraction increase along with the
grain refinement. The addition of refiner is beneficial for forming the equiaxed grain zone. Additionally, refiner
particles dispersed uniformly in the melt causes a large quantity of equiaxed grains formation. The growth of
these nuclei will release a great amount of latent heat, which prohibits their further growth. In addition, the
formation and growth of many equiaxed grains impede the growth of columnar grains.
The decrease of porosity in the specimen with grain refinement is due to the fact that the alloy flow distance is
increased with grain refinement. The fluidity of two different conditions is tested by spiral fluidity. The fluidity is
360 mm at the condition of coarse grain and that of the fine grain is 371 mm. Dahle et al [16]
also reported that
finer grain size should improve fluidity of molten aluminum. This is due to grain refinement postponing dendrite
coherency.
The important consequence of the solidification in superalloy K4169 is the segregation of Nb and the formation of
Laves phase. Laves phase is a brittle inter-metallic topologically close-packed phase with hexagonal structure,
known for its detrimental effect on mechanical properties at room temperature [17]
. The main reason of laves
formation is Nb and Ti segregation [17]
. Figure 6 shows the correlation of grain size and segregation ratios in the
K4169 superalloy. The segregation ratio is defined as the average concentration in the dendritic core over the
average concentration in inter-dendritic region. The segregation ratio close to 1 indicates that the elements can
reduce segregation. It can be clearly seen that the segregation of Al, Mo, Cr and Fe has little change, whereas
the segregation of Nb and Ti decreases with decreasing grain size. So, this is the most important reason for the
decreasing of the quantity of Laves phase.
Fig. 6: Relationship between dendrite segregation ratio and grain size
The increase of the mechanical properties at room temperature in the grain refined samples is mainly due to the
increase of the grain boundaries, which inhibits dislocation slide, and increases the yield strength and ultimate
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strength. At room temperature, the strength of the grain boundary is higher than that of the grain interior [2, 7, 18]
.
Therefore, the crack propagation would be impeded
when encountering a grain boundary. The carbides and Laves phase in fine-grain castings are smaller than those
in the coarse grain, which can also increase the yield and ultimate strength. However, high density and large size
of micro-porosities in the coarse grain samples will lead to the test bars premature fracture, and cause low
elongation and ultimate tensile strength.
4 Conclusions
The effects of the refiners on the as-cast structures and tensile properties of K4169 superalloy were studied. The
results are summarized as follows:
1. When adding mixed refiner of Co3FeNb2 and CrFeNb to the melt of K4169 superalloy, the equiaxed grain
size could be refined and the proportion of equiaxed grains at cross-section could be increased in the
samples with pouring temperature of 1,470- 1,520 °C. Refiner particles with good lattice compatibility with
matrix act as substrata of matrix, thereby causing grain refinement.
2. As the grain refines, the amount of Laves phase decreases and its morphology changes from island to
blocky structure. The carbides in the fine grain samples are fine and dispersive.
3. The amount of porosity in the specimen could be reduced greatly after grain refinement due to the alloy flow
distance being increased with grain refinement.
4. Yield strength and ultimate tensile strength at room temperature increases significantly due to grain
refinement.
When the grain size of K4169 superalloy is 1.85 mm, the highest tensile and yield strength obtained are 1,189.32
MPa and 1,138.47 MPa, respectively.peralloy is 1.85 mm, the highest tensile and yield strength obtained are
1,189.32 MPa and 1,138.47 MPa, respectively.
References
1. Liu Lin, Huang Taiwen, Xiong Yuhua, et al. Grain refinement of superalloy k4169 by addition of refiners: cast
structure and refinement mechanisms. Materials Science and Engineering A, 2005, 394: 1-8.
2. Du Beining, Yang Jinxia, Cui Chuanyong, et al. Effect of grain refinement on the microstructure and tensile
behavior of K417G superalloy. Materials Science and Engineering A, 2015,59-67.
3. Ma Yue, Sun Jiahua, Xie Xishang, et al. An investigation on fine-grain formation and structural character in
cast IN718 superalloy. Journal of Materials Processing Technology, 2003,35-39.
4. Xiong Yuhua, Wei Xiuying, Du Jun, et al. Grain refinement of superalloy IN718C by the addition of
inoculants. Metallurgical and Materials Transaction A, 2004, 35(7): 2111-2114.
5. Liu Lin, Zhang Rong, Wang Liuding, et al. A new method of fine grained casting for nickle-base superalloys.
Journal of Materials Processing Technology, 1998, 77: 300-304.
6. Ma Xiaoping, Li Yingju and Yang Yuansheng. Grain refinement effect of pulsed magnetic field on solidified
microstructure of superalloy IN718. Journal of Materials Research, 2009, 24(10): 3174-3181.
7. Wei C N, Bor H Y, Ma C Y, et al. A study of IN713LC superalloy grain refinement effects on microstructure
and tensile properties. Materials Chemistry and Physics, 2003, 80(1): 89-93.
8. Jin Wenzhong, Bai Fudong, Li Tingju, et al. Grain refinement of superalloy IN100 under the action of rotary
magnetic fields and inoculants. Materials Letters, 2008, 62(10-11): 1585-1588.
9. Bashir S and Thomas M C. Effect of interstitial content on High-temperature fatigue crack propagation and
Low-cycle fatigue of alloy 720. Journal of Materials Engineering and Performance, 1993, 2(4): 545-550.
10. Miao Jiashi, Pollock T M and Jones J W. Crystallographic fatigue crack initiation in Nickel-based superalloy
Rene′ 88DT at elevated temperature. Acta Materialia, 2009, 57(20): 5964-5974.
11. Liu R, Xi S Q, Kapoor S, et al. Effect of chemical composition on Solidification, microstructure and hardness
of Co-Cr-W-Ni and Co-Cr-Mo-Ni alloy systems. IJRRAS, 2010, 5(2): 110-122.
12. Li Ailan, Tang Xin, Gai Qidong, et al. Effect of heat treatment on microstructure of K4169 superalloy. Journal
of Aeronautical Materials, 2006, 26(3): 311-312. (In Chinese)
BULETIN - APLINDO No.48/2016
48
13. Li Yamin, Liu Hongjun, Liu Jie, et al. Effect of Zr addition on precipitates in K4169 superalloy. China
Foundry, 2012, 9(1): 6-10.
14. Bramfitt B L. The effect of carbide and nitride additions on the heterogeneous nucleation behavior of liquid