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Refinement Mechanism Research of Al3Ni Phase in Ni-7050 Alloy
Rare Metal Materials and Engineering Volume 42, Issue 1, January 2013 Online English edition of the Chinese language journal Cite this article as: Rare Metal Materials and Engineering, 2013, 42(1): 0006-0013.
ARTICLE
Refinement Mechanism Research of Al3Ni Phase in Ni-7050 Alloy Gao Peng1, 1
Zhou Tietao1,
Xu Xiaoqing2,
Gao Zhi1,
Chen Li1
2
Beihang University, Beijing 100191, China; Inner Mongolia Normal University, Huhhot 010022, China
Abstract: This work investigated the refinement of Al3Ni phase in Ni-7050 composite, which was prepared by the melting reaction with two different components, 5% and 10% Ni added in 7050 aluminium alloy. Hardness tests, metallographic observation, SEM and DSC were used to analyze the structures and performances. The results show that the hardness of the two composites was obviously increased to 1918 and 2364 MPa after T6 aging. Fracture analysis shows that the fracture mechanism was mainly Al3Ni brittle fracture and interfaces debond. Experiments were designed to investigate the influences of annealing/rolling treatment on structures and properties of the composites. The results indicated: (1) Long time annealing treatment could induce refinement and spheroidization of the Al3Ni phase. With the annealing time prolonged, the hardness first increased and then decreased. (2) After multi-step hot/cold rolling, the microscopic structure of composites showed that the size of Al3Ni phase changed progressively, Al3Ni phase was smashed and transformed from original long slab to nearly isometric particles. Key words: aluminium alloy; Al3Ni; annealing; rolling; hardening
The materials with high strength, high wear resistance, high elongation, low density and good high-temperature performance never failed in the fascination of metallurgical and materials scientists. 7000 series aluminum alloy is a kind of alloy with high strength and high toughness, widely used in the aerospace industry. At present time, the yield strength of ultra-high strength aluminium alloy is about 500~650 MPa. According to Mackenzie J.K.’s calculation about FCC structure using the solid mechanical method, we know that τmax≈G/30, and the shearing elastic modulus of aluminium alloy is about 26 GPa. The maximum strength can reach up to 870 MPa in theory. And on the basis of Von Mises yield rule of dislocation theory, the tensile yield strength can reach up to 1148 MPa. It is obvious to see that the aluminium alloy’s properties still can achieve a large advancement. For strengthening aluminums alloy, there are several methods used widely, such as solution, precipitation, refine-grain, dispersion strengthening and compound strengthening. The most resultful and important method is precipitation strengthening. The increasing of precipitate’s
volume fraction could improve the yield strength effectively. When the volume fractions reach to 0.1, the strength will be 1000 MPa in theory, but now it is only about 0.02 or 0.03. Second, the refine-grain strengthening is also effective to the alloy. According to some research results, when the grain size arrived at 1 μm, the yield strength can reach to 850 MPa. So the main problem of strengthening aluminium alloy is to improve the precipitate’s volume fraction, to control the size class of precipitate and its best special distribution, and microstructure of precipitate phase. The current research on Al3Ni strengthening aluminium matrix composite material is mainly about Al-Ni binary system [1-3] or adding Ni into 2000 series aluminium alloy [4]. Until now, there’s no research on Ni with 7000 series aluminium alloy. There’re several kinds of Al-Ni intermetallic compounds, with good strength, hardness performance and also excellent stability at high temperature. Also, adding nickel to aluminium matrix can generate in-situ Al3Ni phase, with the phase and matrix performed by metallurgical bonding with elastic modulus at 116~140 GPa
Received date: January 4, 2012 Corresponding author: Zhou Tietao, Professor, School of Material Science and Technology, Beihang University, Beijing 100191, P. R. China, Tel: 0086-10-82317125, E-mail: [email protected] Copyright © 2013, Northwest Institute for Nonferrous Metal Research. Published by Elsevier BV. All rights reserved.
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Gao Peng et al. / Rare Metal Materials and Engineering, 2013, 42(1):0006-0013
and melting at about 890 ℃ [5,6]. The aim of this paper is to improve the precipitate strengthening ability with nickel addition, increase the strength and hardness and to find a way to control the precipitated phase.
2.1
Table 2
Zr 0.1 0.1
Ti 0.02 0.02
Mn 0.01 0.01
Ni 5 10
Result of DSC analysis
a
b
25 µm Fig.1
Casting structure of material A (a) and material B (b) a
b
Spectrum 2 Spectrum 1
20 µm
Results of DSC tests
Casting structure
Seen from Fig.1, the cast structures of materials A and B are clearly different. When the content of Ni reaches to 5%, it is closer to the eutectic point, but because the existence of Zn, Mg, Cu and other elements, the actual composition is more than the eutectic point; therefore we could see some pro-eutectic phases besides a large number of dendritic eutectic microstructure in the material A(Fig.1a). As the content of Ni in material B is far away from the eutectic point, we could see lots of coarse pro-eutectic phase in Fig.1b, while the eutectic organization is particularly too small to recognize. The result of the energy spectrum analysis on the pro-eutectic phase is shown in Fig.2 and Table 3. In the marginal area (Fig.2a), the atomic percentages of Al and Ni are 61.64% and 21.16%, with Cu (15.03%), Zn (1.58%) and Fe (0.595). In the center (Fig.2b), the
Al Bal. Bal.
Temp. of start-melting Temp. of fully-melting Temp. of peak /ºC /ºC /ºC A 473.2 488.1 478.4 B 473.1 487.7 478.1
Results and Analysis
The results of DSC analysis were given in Table 2. According to the results of DSC analysis, we found that the two materials show very similar melting behavior. So we chose 470 ºC as the temperature of solution and annealing treatment. The solution and aging treatment was applied as follows: 470 ºC (0.5 h) + water quenching +120 ºC (24 h).
2.2
fraction, %) Mg Cu 2.09 2.18 1.98 2.07
Experimental
The matrix material was 7050 aluminium alloy, whose composition was mainly 6.1% Zn, 2.2% Mg, 2.3% Cu, 0.11% Zr, 0.02% Ti, 0.01% Mn (all in wt%). And Ni was added in two different ratios denoted as A (5 wt %-Ni) and B (10wt% -Ni). We melted the materials in a vacuum melting furnace, and the nominal compositions of the studied material are given in Table 1. The melting point was measured by DSC, to assure its solution and annealing temperature. Then annealing treatment was used to improve the precipitated phase. T6 aging treatment (120 ºC, 24 h) and rolling (room-temperature and 450 ºC) were also used to refine the precipitated phase. Fracture and component analysis were completed on the scanning electron microscope with an energy spectrum. The polished specimens were observed by optical microscope after corrosion. Hardness curve was tested after solution, annealing and the aging treatment.
2
Nominal compositions of the two materials (mass
Zn 5.80 5.49
A B
c Intensity/a.u.
1
Table 1
d
Al
Al Zn Ni Cu Fe
0
NiCuCu FeFe NiZn Zn
2
4
6
8
Energy/keV Fig.2
10
Ni Zn
0
Ni NiZn Zn
2
4
6
8
10
Energy/keV
SEM images (a, b) and the corresponding energy spectra (c, d) of Al-Ni phase: (a, c) marginal area and (b, d) centre area
atomic percentages of Al and Ni are 73.42% and 25.91%, with Zn (0.67%). So from Table 3 we can confirm that the phase is mainly composed of Al and N, which can be known as Al3Ni. And XRD test results also prove that there are Al3Ni phase (Fig.3). We could easily find that diffraction peak of Al3Ni in material B is higher than that in A.
2.3
Fracture analysis
To study the fracture properties, the tensile specimens from the two materials were prepared, and the fracture surface was shown in Fig.4.
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Gao Peng et al. / Rare Metal Materials and Engineering, 2013, 42(1):0006-0013
Table 3
Results of energy spectrum analysis for this first phase in Fig.2a and 2b Spectrum 1
Element
ω/% 41.62 31.08 2.58 23.89 0.83 100.00
Al Ni Zn Cu Fe Total
Spectrum 2
at/% 61.64 21.16 1.58 15.03 0.59 100.00
ω/% 55.88 42.92 1.20 100.00
at/% 73.42 25.91 0.67 100.00
1400
a
Intensity/cps
1200
Al Al3 Ni MgZn2
1000 800 600 400 200 0 20
30
40
50
Intensity/cps
8000
60
70
80
90 b
Al Al3 Ni MgZn2
6000
100
4000 2000
3
0 30
40
50
60
70
80
90
100
3.1.1
2θ/(º)
XRD patterns: (a) material A and (b) material B a
b
100 µm
Fig.4
Methods to Refine the Precipitations
3.1 20
Fig.3
expand along the interface after bypassing the phase, some large size Al3 Ni phases fractured under tensile stress (the white arrows) or debonded with the matrix at the interface (the black arrows). From this, the fracture mechanism of this material is mainly Al 3 Ni fracture or interface debonding. The elastic modulus of Al3Ni phase is 1.5 times higher than that of aluminum. Under the action of external tensile stress, the stress imposed on the Al3Ni phase is much larger than that on the matrix alloy. With the larger size of the reinforcement, it is more possible that the withstanding stress exceeds its tensile strength, so the fracture is controlled by the reinforcement [7]. Thus, it might lead to strength reduction of the high-strength aluminum body after the entry of Al3Ni reinforcement; in addition, the presence of defects such as micro-cracks in the Al3Ni phase may further reduce the fracture strength; when the relative size of the Al3Ni phase was larger, the probability of the existence of defects was higher. There is high shear stress in the Al/Al3Ni interface; which would rise rapidly when there is a load tensile force. When the shear stress exceeded the bond strength, it would lead to interface debonding in Al/Al3Ni interface, and then cause cracks. Therefore, to improve the interfacial bond strength of Al matrix/Al3Ni or to change the morphology and distribution of the phase by heat treatment, plastic processing and other methods were expected to improve the mechanical properties of this material.
Fracture surface of casting tensile samples: (a) material A and (b)material B
There were marked fracture traces of Al 3 Ni with big size and cracks in the tensile fracture surface as shown in Fig.4. In the crack propagation process, it was obviously hindered by the Al3 Ni phase. When cracks extended close to Al 3 Ni phase, its direction changed and continued to
Annealing treatment and research Annealing treatment
To study the impact on the morphology and distribution of the Al3Ni phase, we observed annealed specimens with different time. And the specimens with annealed + T6 aging treatment have similar condition. Seen from the metallographic photos (Fig.5), both in material A and B, almost all Al3Ni plates were fused and spheroidized. The Al3Ni plates were refined obviously after 75 h annealing treatment (or annealing + aging). And all the plates dispersed more uniformly in the matrix as annealing time increased, with the average size about 3 μm in material A. For there were lots of thick pro-eutectic phases in material B, even after annealing + aging treatment the average size was about 8 μm. Comparing annealing with and without aging treatment, it is found that Al3Ni phase in material A and B both became dispersed. So it could be illustrated that annealing treatment can decrease aliquation and agglomeration of the dendritic eutectic structure. 3.1.2 Al3Ni phase analysis Using metallographic analysis methods, we calculated the Al3Ni plates’ average size class of material A, and B in
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Gao Peng et al. / Rare Metal Materials and Engineering, 2013, 42(1):0006-0013
Table 4 a
Average size of Al3Ni phase in different annealing time (μm)
b Material
0h
24 h 48 h 75 h
100 h
150 h
200 h
A
11.5
10.6
8.6
5.8
3.9
3.2
3.1
B
18.5 16.9 13.3
11.6
9.8
9.2
8.4
b a c
d
Al3Ni phase
25 µm
Fig.5
OM images of material A (a, b), and B (c, d) with annealing treatment: (a,c) 24 h, and (b, d) 200 h
2σ Vm r
μ = G m + (1 − x )
∂G m ∂x
Fig.6
Simplified model of Al3Ni phase before fusing and spheroidizing
different annealing time (for every different status we chose more than ten photos to statistic, and the plates’ statistic result showed the average length of major axis and minor axis), shown in Table 4. The average size of Al3Ni plates became smaller as the annealing time extended. When the annealing time extended to 200 h, the average size of Al3Ni plates can reach to 3.1 μm in material A and 8.4 μm in material B. This could illustrate long time annealing treatment can decrease the average size of Al3Ni plates clearly. And during 24 to 100 h, the effects of refining Al3Ni plates are more significant. This could be found both in the specimens with annealing and with annealing + aging. 3.1.3 Spheroidizing mechanism analysis In order to explain the reasons for the large Al3Ni plates fusing and spheroidizing to small plates after annealing, we built the model in Fig.6. Fig.6 is the simplified diagram of Al3Ni plates before fusing and spheroidizing. The curvature radius of point a, b is ra and rb. Obviously, ra is much less than rb. By the thermodynamic formula [8] G m = G m (0) +
Matrix
(1) (2)
Gm (0): The molar Gibbs free energy when the curvature radius is zero r: The curvature radius σ: The surface tension Vm: The molar volume X: The mole fraction of group element The chemical potential difference of two points was:
∂Gm(a) ∂Gm (b) − (1 − Xb ) ∂Xa ∂Xb (3) ∂Gm (a) ∂Gm (b) 2σ Vm 2σ Vm = − + (1 − X a ) − (1 − X b ) ra ∂X a ∂X b rb
Δμ = Gm(a) − Gm (b) + (1 − Xa )
Because Xa = Xb
Δμ =
2σ V m 2σ V m − ra rb
(4)
Since ra is much less than rb, the chemical potential at point a is higher than that at point b in the matrix. According to Ref.[9], at 400 ºC, the Ni atoms have been able to spread long distance, so it is enough for Ni atoms to carry out a long-ranged diffusion at 470 ºC. The Ni atoms would diffuse from point a with higher chemical potential to the point b with lower chemical potential[10], which brought damage to the equilibrium concentration of Ni atoms near the point a. In order to keep balance, the Al3Ni dissolved continuously to make up the shortage of nickel concentration, so the groove would be deepening and widening continuously, the sharp corner would be dissolving constantly, until the branch had fused, and the sharp corner had disappeared and spheroidized. The fusing of the branch is bound to create new grooves and sharp corners, and then it would continue this process. After repeated fusing of branches and spheroidizing of sharp corners, we eventually obtained tiny Al3Ni grains which were nearly spheroidal, see in Fig.5b.
3.1.4
Hardness test
We tested the hardness after annealing and T6 aging treatment (the load was 1 kg, and load-on 30 s. Average value from ten specimens), the results were shown in Table 5 and Fig.7.
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Gao Peng et al. / Rare Metal Materials and Engineering, 2013, 42(1):0006-0013
Table 5
Results of Vickers hardness test (HV) after annealing(MPa)
Annealing time
Sample A
Sample B
7050
0 h+T6
1918
2364
180
24 h+T6
1980
2385
—
48 h+T6
2050
2385
—
75 h+T6
2080
2395
—
100 h+T6
1957
2293
—
150 h+T6
1952
2278
—
200 h+T6
1950
2272
—
Hardness, HV/MPa
2400
Material A Material B
2300 2200 2100 2000 1900
0
50
100
150
200
Annealing Time/h Fig.7
Curves of hardness-annealing time
phase is the main reason of fracture. So after annealing treatment experiments, we tried to use traditional deformation method to lower the size of the large pro-eutectic phase and reduce its brittleness. In following experiments, Cold-Rolling and Hot-Rolling were adopted to smash the Al3Ni plates. The original thickness of sample was 10 mm, every time the rolling reduction is 1 mm, and after seven-step rolling, the total deformation rate was 70%. The microstructure of Al3Ni phases in different deformation is shown in Fig.8. (H-R stands for Hot-Rolling and C-R stands for Cold-Rolling in the following.) The microstructure showed that, when the C-R deformation rate was 10%, a few large Al3Ni plates were fractured. Al3Ni long plates which were parallel to the rolling stress need large deformation and the possibility of fracture is big, while the long plates which were perpendicular to the rolling stress can have compressive deformation to offset the deformation. Most of the Al3Ni plates were longer than 50 μm, the average space between Al3Ni plates was large, about 40 μm, and some Al3Ni plates had cracks in their cusps. The H-R showed similar situation to C-R, but most of the Al3Ni plates were integrated, cracks can be easily found but clear fractures had not existed. Because at the high temperature (450 ºC), Al-matrix was ductile and easy to deform, the Al-matrix had compressive deformation which was the main effect of the compressive stress, and Al3Ni plates mainly had position change.
The hardness of the aluminum matrix composite alloy were much higher than that of 7050 alloy, and increased with the content of Ni, which indicated that the hardness of 7050 alloy could be significantly improved by Al3Ni phases. From the Fig.7, with the prolonging of annealing time, the hardness first increased and then decreased, especially for the material A who has near eutectic composition. This phenomenon may be caused by the following reasons: (1) the casting defects of the materials (such as porosity, segregation, dispersed shrinkage, shrinkage cavity and so on) were eliminated by the annealing treatment, which made the structure more compact; (2) the Al3Ni is hard and brittle phase, so their presence might significantly improve the hardness of materials, especially when they were dispersed in Al matrix, and it is particularly reflected on the material A, whose hardness was improved much more than that of the material B; (3) as the annealing time was further extended, the dispersion effect of Al3Ni phase on the increase of hardness would not be strengthened, but the matrix would have the softening trend because of the elimination of residual stress, which led to the decline in hardness. 3.2 Rolling treatment and research 3.2.1
Cold-Rolling/Hot-Rolling
From the former part, we know the break of large Al3Ni
a
b
c
d
e
f
50 µm
Fig.8
Microstructure of Al3Ni after H-R/C-R: (a) 10%C-R, (b) 40%C-R, (c) 70%C-R, (d) 10%H-R, (e) 40%H-R, and (f) 70%H-R)
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Gao Peng et al. / Rare Metal Materials and Engineering, 2013, 42(1):0006-0013
Average Length/µm
11
a
10
Minor axis of C-R Minor axis of H-R
9 8 7 6 5 4 3 10
20
30
40
50
60
120
70 b
Major axis of C-R Major axis of H-R
100 Average Length/µm
During deformation rate change from 20% to 40%, the Cold-Rolling and Hot-Rolling showed the same trend, the Al3Ni plates became longer continuously. When the deformation was about 40%, the Al3Ni plates reach the extreme of length direction extension, and long Al3Ni plates were inclined to rupture owing to the shearing stress. When the deformation rate is 50%, long Al3Ni plates were smashed into short plate of which the major axis to minor axis ratio is about 5:1; at that time, most of the long Al3Ni plates had disappeared, breakage effect of the compressive stress began to diminish but the compressive effect of Al3Ni short plate began to increase. During the deformation rate change from 60% to 70%, the short Al3Ni plates attenuated and became secondary long Al3Ni plate, then obvious rupture appeared again, and finally the Al3Ni plates were crushed into grains of which the major axis to minor axis ratio is about 2:1. Different from cold-rolling, Al3Ni grains of 70% deformation after hot-rolling had clearer cusps than that after cold-rolling. It is because during the hot rolling, the yield strength of Al-matrix is low and Al-matrix was ductile, which transmit less stress to Al3Ni phase than that under cold-rolling, so the crushing effect was less obvious. Under the cold-rolling condition, Al3Ni plates can squeeze each other to eliminate the cusps, but under the hot-rolling condition, the stress between Al3Ni plates could be transmitted to Al-matrix and cause Al-matrix deformation, so cusps were preserved finally. 3.2.2 Al3Ni phase analysis Since the Al3Ni phase shape is mainly plate, we made the hypothesis that the cross-section of Al3Ni phase can be considered as symmetrical rectangle, four photos of each rolling state were chosen for data analysis and 128 Al3Ni plates were measured in the four photos, measuring objects included the lengths of major axis and minor axis of Al3Ni plates. And the curve of the Al3Ni plates’ average length in different deformation rates was made and is shown in Fig.9. The curves in Fig.9a show that both the H-R and C-R can reduce the length of minor axis. After 70% deformation of H-R, the average minor axis is about 4.4 μm comparing to 3.4 μm of C-R. The curves also indicated that, with the increase of deformation rate under C-R, the length of minor axis of Al3Ni phase decreased continuously, but when the deformation rate is 60% under H-R, the minor axis did not attenuate obviously after 7th hot rolling and maintained 4.4 μm. during the 7th hot rolling, the main effect of compressive stress was to smash the Al-Ni plate further, and the average length of major axis was reduced from 37.5 to 25.8 μm Since the influence of rolling on the minor axis of Al3Ni phase is inclined to be linear, curves were fitted as below(X represents deformation rate and Y represents the length of minor axis).
80 60 40 20 0 10
20
30
40
50
60
70
Deformation Rate/% Fig.9
Curves of the Al3Ni plates’ average length in different deformation rates of cold-rolling (C-R) and hot-rolling (H-R)
YC-R=–11.89X+ 11.51 YH-R=–7.607X + 9.214 It can be found that the reducing effect of minor axis length of cold-rolling is 56% better than that of hot-rolling. Fig.9b indicates that there exist peak value of the major axis length during the multi-step hot rolling process. When the deformation rate is 30%, the average length reached the maximum of 106 μm. From Fig.8, we know the Al 3Ni plates surfaces were full of cracks but not completely fractured, but when the deformation rate reached 40%, cracks were enlarged and one long plate became several short plates, and the average length of major axis decreased to 80 μm. From the OM photos, fractures mainly took place on long Al3Ni plates; short Al3Ni plates attenuated but did not fracture. Different from the H-R, there were two peak values of major axis length during the multi-step C-R process; they were 105 μm of 30% deformation and 37.6 μm of 60% deformation. After 30% deformation, there were many cracks which inclined to the major axis at an angle of 45º but did not fracture completely. After 40% deformation, the average major axis length suddenly dropped to 40 μm, most the long Al3Ni plates had been smashed into secondary Al3Ni plate, and the major axis to minor axis ratio is about
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6.4:1, close to that of the 10% deformation condition. When the deformation rate is 60% which, another peak value of major axis length appeared, was 37.6 μm. It is because the focus of compressive stress in Al3Ni plates is the determinants. 3.2.3 Hardness test We tested the hardness after every step of cold/hot-rolling treatment (the load was 1 kg, and load-on 30 s), and each specimen tested ten points and calculated the average value. The results were shown in Table 6. Then we studied the hardness curve of Al-matrix (average value from ten specimens) in Cold-Rolling and Hot-Rolling process, with microhardness instrument to avoid the Al3Ni phases (the load was 0.0 1kg, load-on 10 s). showed in Table 7. As the deformation rate increased, the hardness of material B increased obviously both in Cold-Rolling or Hot-Rolling process. And at 30%, 60% deformation rate, their hardening effect is not such obviously. In Cold-Rolling process, work-hardening effect is the main reason of hardness increased. And in Hot-Rolling process, the hardness increased and with a peak value. After the 30% deformation rate, the soften effect was larger. So the hardness of Al-matrix had fallen down. The hardness of material B increased with deformation rate. But at 60% (C-R) and 30% (H-R) deformation rate, there was a minimum value, and it was corresponding with the change of Al3Ni plates’ size. In the former condition, the Al3Ni plates became larger, so the precipitation hardening effect was significantly weakened by the size effect. Table 6
Results of Al-Ni phases’ Vickers hardness test (HV) after C-R/H-R(MPa) 10%
20%
30%
40%
50%
60%
70%
C-R
921
969
974
992
1041
1015
1090
H-R
925
932
947
935
937
955
1032
Table 7
Results of Vickers hardness (HV) in Al matrix after multi-step C-R/H-R(MPa) 10%
4
20%
30%
40%
50%
60%
70%
C-R
686
705
708
714
757
776
807
H-R
682
708
740
722
695
674
661
Discussion and Conclusions
1) From Ref.[13,14], the effect on mechanical property of the material after Al3Ni phase generated could be illustrated as the following formula: First, the effect of Al3Ni plates on strength can be expressed as σ y = σ mVm +σ r Vr r
(5)
σy: The calculated valued of yield strength σm, σr: The yield strength of matrix and reinforcement particle Vm, Vr: The volume fraction of matrix and reinforcement particle And actually, the strength of material is not only affected by the volume fraction. When the size of reinforcement particle is large enough, the effect of Orowan bypassing dislocation strengthening mechanism would be ignored. In the external stress, the slipping of dislocations in matrix was slowed down at interface of matrix/Al3Ni plates, with stress concentration in the Al3Ni plates. The stress value σy could be described as [14]:
σy=
3G m G p b V p 2 d (1−V p )c
(6)
Gm, Gp: the shear modulus of matrix and Al3Ni plate b: Burgers vector Vp: the volume fraction of Al3Ni plate c: constant d: average size of Al3Ni plates By substituted relevant parameters, we can infer that when the average size of Al3Ni phase reaches to 5 μm and the volume fraction of Ni reaches to 30%Al3Ni, the yield strength of the material can get to 630 MPa, far more than 500MPa of 7050 aluminum alloy. 2) Through the fracture analysis, the fracture mechanism of the material is mainly Al3Ni fracture and interface debonding. And the plasticity decreased as the content of Ni increased. 3) As the annealing time extending, the Al3Ni phase had a trend of fusing and spheroidizing. After long time annealing, the aluminum matrix composites could obtain tiny and dispersed Al3Ni grains which were nearly spheroidal. The average diameter of the Al3Ni grains is about 3 μm in the material A (5%Ni), and it is 6 μm in the material B (10%Ni). 4) Besides annealing, multi-step Hot-rolling/Cold-rolling process also can make the Al3Ni plates size decrease obviously. The minor axis length of Al3Ni plates decreased continuously due to the compressive effect and smashing effect, and the length in minor axis at 70% deformation rate is about 4 μm. Comparing to the minor axis, there existed peak values in major axis length, mainly because of the compressive effect which squeeze the Al3Ni plates to be longer before the stress reached to the threshold value. When the stress reached to the threshold value, the Al3Ni plates fractured and then the major axis length dropped rapidly. From statistical results, the effect of Cold-rolling was slightly better than that of hot-rolling. 5) The addition of Ni can significantly improve the hardness of 7050 aluminum alloy. After T6 aging, the
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Gao Peng et al. / Rare Metal Materials and Engineering, 2013, 42(1):0006-0013
hardness of material A (5%Ni) and material B (10%Ni) were 1918 and 2364 MPa respectively, and both were much higher than 180 HV, the hardness of the 7050 deformed aluminum alloy. As the annealing time extending, the hardness of all first increased and then decreased, especially for material A (5%Ni), whose hardness could be up to 2070 MPa after 48 h annealing and T6 aging, and the hardness of material B (10%Ni) could be up to 2385 MPa after 48 h annealing and T6 aging.
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ARTICLE
Refinement Mechanism Research of Al3Ni Phase in Ni-7050 Alloy Gao Peng1, 1
Zhou Tietao1,
Xu Xiaoqing2,
Gao Zhi1,
Chen Li1
2
Beihang University, Beijing 100191, China; Inner Mongolia Normal University, Huhhot 010022, China
Abstract: This work investigated the refinement of Al3Ni phase in Ni-7050 composite, which was prepared by the melting reaction with two different components, 5% and 10% Ni added in 7050 aluminium alloy. Hardness tests, metallographic observation, SEM and DSC were used to analyze the structures and performances. The results show that the hardness of the two composites was obviously increased to 1918 and 2364 MPa after T6 aging. Fracture analysis shows that the fracture mechanism was mainly Al3Ni brittle fracture and interfaces debond. Experiments were designed to investigate the influences of annealing/rolling treatment on structures and properties of the composites. The results indicated: (1) Long time annealing treatment could induce refinement and spheroidization of the Al3Ni phase. With the annealing time prolonged, the hardness first increased and then decreased. (2) After multi-step hot/cold rolling, the microscopic structure of composites showed that the size of Al3Ni phase changed progressively, Al3Ni phase was smashed and transformed from original long slab to nearly isometric particles. Key words: aluminium alloy; Al3Ni; annealing; rolling; hardening
The materials with high strength, high wear resistance, high elongation, low density and good high-temperature performance never failed in the fascination of metallurgical and materials scientists. 7000 series aluminum alloy is a kind of alloy with high strength and high toughness, widely used in the aerospace industry. At present time, the yield strength of ultra-high strength aluminium alloy is about 500~650 MPa. According to Mackenzie J.K.’s calculation about FCC structure using the solid mechanical method, we know that τmax≈G/30, and the shearing elastic modulus of aluminium alloy is about 26 GPa. The maximum strength can reach up to 870 MPa in theory. And on the basis of Von Mises yield rule of dislocation theory, the tensile yield strength can reach up to 1148 MPa. It is obvious to see that the aluminium alloy’s properties still can achieve a large advancement. For strengthening aluminums alloy, there are several methods used widely, such as solution, precipitation, refine-grain, dispersion strengthening and compound strengthening. The most resultful and important method is precipitation strengthening. The increasing of precipitate’s
volume fraction could improve the yield strength effectively. When the volume fractions reach to 0.1, the strength will be 1000 MPa in theory, but now it is only about 0.02 or 0.03. Second, the refine-grain strengthening is also effective to the alloy. According to some research results, when the grain size arrived at 1 μm, the yield strength can reach to 850 MPa. So the main problem of strengthening aluminium alloy is to improve the precipitate’s volume fraction, to control the size class of precipitate and its best special distribution, and microstructure of precipitate phase. The current research on Al3Ni strengthening aluminium matrix composite material is mainly about Al-Ni binary system [1-3] or adding Ni into 2000 series aluminium alloy [4]. Until now, there’s no research on Ni with 7000 series aluminium alloy. There’re several kinds of Al-Ni intermetallic compounds, with good strength, hardness performance and also excellent stability at high temperature. Also, adding nickel to aluminium matrix can generate in-situ Al3Ni phase, with the phase and matrix performed by metallurgical bonding with elastic modulus at 116~140 GPa
Received date: January 4, 2012 Corresponding author: Zhou Tietao, Professor, School of Material Science and Technology, Beihang University, Beijing 100191, P. R. China, Tel: 0086-10-82317125, E-mail: [email protected] Copyright © 2013, Northwest Institute for Nonferrous Metal Research. Published by Elsevier BV. All rights reserved.
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Gao Peng et al. / Rare Metal Materials and Engineering, 2013, 42(1):0006-0013
and melting at about 890 ℃ [5,6]. The aim of this paper is to improve the precipitate strengthening ability with nickel addition, increase the strength and hardness and to find a way to control the precipitated phase.
2.1
Table 2
Zr 0.1 0.1
Ti 0.02 0.02
Mn 0.01 0.01
Ni 5 10
Result of DSC analysis
a
b
25 µm Fig.1
Casting structure of material A (a) and material B (b) a
b
Spectrum 2 Spectrum 1
20 µm
Results of DSC tests
Casting structure
Seen from Fig.1, the cast structures of materials A and B are clearly different. When the content of Ni reaches to 5%, it is closer to the eutectic point, but because the existence of Zn, Mg, Cu and other elements, the actual composition is more than the eutectic point; therefore we could see some pro-eutectic phases besides a large number of dendritic eutectic microstructure in the material A(Fig.1a). As the content of Ni in material B is far away from the eutectic point, we could see lots of coarse pro-eutectic phase in Fig.1b, while the eutectic organization is particularly too small to recognize. The result of the energy spectrum analysis on the pro-eutectic phase is shown in Fig.2 and Table 3. In the marginal area (Fig.2a), the atomic percentages of Al and Ni are 61.64% and 21.16%, with Cu (15.03%), Zn (1.58%) and Fe (0.595). In the center (Fig.2b), the
Al Bal. Bal.
Temp. of start-melting Temp. of fully-melting Temp. of peak /ºC /ºC /ºC A 473.2 488.1 478.4 B 473.1 487.7 478.1
Results and Analysis
The results of DSC analysis were given in Table 2. According to the results of DSC analysis, we found that the two materials show very similar melting behavior. So we chose 470 ºC as the temperature of solution and annealing treatment. The solution and aging treatment was applied as follows: 470 ºC (0.5 h) + water quenching +120 ºC (24 h).
2.2
fraction, %) Mg Cu 2.09 2.18 1.98 2.07
Experimental
The matrix material was 7050 aluminium alloy, whose composition was mainly 6.1% Zn, 2.2% Mg, 2.3% Cu, 0.11% Zr, 0.02% Ti, 0.01% Mn (all in wt%). And Ni was added in two different ratios denoted as A (5 wt %-Ni) and B (10wt% -Ni). We melted the materials in a vacuum melting furnace, and the nominal compositions of the studied material are given in Table 1. The melting point was measured by DSC, to assure its solution and annealing temperature. Then annealing treatment was used to improve the precipitated phase. T6 aging treatment (120 ºC, 24 h) and rolling (room-temperature and 450 ºC) were also used to refine the precipitated phase. Fracture and component analysis were completed on the scanning electron microscope with an energy spectrum. The polished specimens were observed by optical microscope after corrosion. Hardness curve was tested after solution, annealing and the aging treatment.
2
Nominal compositions of the two materials (mass
Zn 5.80 5.49
A B
c Intensity/a.u.
1
Table 1
d
Al
Al Zn Ni Cu Fe
0
NiCuCu FeFe NiZn Zn
2
4
6
8
Energy/keV Fig.2
10
Ni Zn
0
Ni NiZn Zn
2
4
6
8
10
Energy/keV
SEM images (a, b) and the corresponding energy spectra (c, d) of Al-Ni phase: (a, c) marginal area and (b, d) centre area
atomic percentages of Al and Ni are 73.42% and 25.91%, with Zn (0.67%). So from Table 3 we can confirm that the phase is mainly composed of Al and N, which can be known as Al3Ni. And XRD test results also prove that there are Al3Ni phase (Fig.3). We could easily find that diffraction peak of Al3Ni in material B is higher than that in A.
2.3
Fracture analysis
To study the fracture properties, the tensile specimens from the two materials were prepared, and the fracture surface was shown in Fig.4.
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Gao Peng et al. / Rare Metal Materials and Engineering, 2013, 42(1):0006-0013
Table 3
Results of energy spectrum analysis for this first phase in Fig.2a and 2b Spectrum 1
Element
ω/% 41.62 31.08 2.58 23.89 0.83 100.00
Al Ni Zn Cu Fe Total
Spectrum 2
at/% 61.64 21.16 1.58 15.03 0.59 100.00
ω/% 55.88 42.92 1.20 100.00
at/% 73.42 25.91 0.67 100.00
1400
a
Intensity/cps
1200
Al Al3 Ni MgZn2
1000 800 600 400 200 0 20
30
40
50
Intensity/cps
8000
60
70
80
90 b
Al Al3 Ni MgZn2
6000
100
4000 2000
3
0 30
40
50
60
70
80
90
100
3.1.1
2θ/(º)
XRD patterns: (a) material A and (b) material B a
b
100 µm
Fig.4
Methods to Refine the Precipitations
3.1 20
Fig.3
expand along the interface after bypassing the phase, some large size Al3 Ni phases fractured under tensile stress (the white arrows) or debonded with the matrix at the interface (the black arrows). From this, the fracture mechanism of this material is mainly Al 3 Ni fracture or interface debonding. The elastic modulus of Al3Ni phase is 1.5 times higher than that of aluminum. Under the action of external tensile stress, the stress imposed on the Al3Ni phase is much larger than that on the matrix alloy. With the larger size of the reinforcement, it is more possible that the withstanding stress exceeds its tensile strength, so the fracture is controlled by the reinforcement [7]. Thus, it might lead to strength reduction of the high-strength aluminum body after the entry of Al3Ni reinforcement; in addition, the presence of defects such as micro-cracks in the Al3Ni phase may further reduce the fracture strength; when the relative size of the Al3Ni phase was larger, the probability of the existence of defects was higher. There is high shear stress in the Al/Al3Ni interface; which would rise rapidly when there is a load tensile force. When the shear stress exceeded the bond strength, it would lead to interface debonding in Al/Al3Ni interface, and then cause cracks. Therefore, to improve the interfacial bond strength of Al matrix/Al3Ni or to change the morphology and distribution of the phase by heat treatment, plastic processing and other methods were expected to improve the mechanical properties of this material.
Fracture surface of casting tensile samples: (a) material A and (b)material B
There were marked fracture traces of Al 3 Ni with big size and cracks in the tensile fracture surface as shown in Fig.4. In the crack propagation process, it was obviously hindered by the Al3 Ni phase. When cracks extended close to Al 3 Ni phase, its direction changed and continued to
Annealing treatment and research Annealing treatment
To study the impact on the morphology and distribution of the Al3Ni phase, we observed annealed specimens with different time. And the specimens with annealed + T6 aging treatment have similar condition. Seen from the metallographic photos (Fig.5), both in material A and B, almost all Al3Ni plates were fused and spheroidized. The Al3Ni plates were refined obviously after 75 h annealing treatment (or annealing + aging). And all the plates dispersed more uniformly in the matrix as annealing time increased, with the average size about 3 μm in material A. For there were lots of thick pro-eutectic phases in material B, even after annealing + aging treatment the average size was about 8 μm. Comparing annealing with and without aging treatment, it is found that Al3Ni phase in material A and B both became dispersed. So it could be illustrated that annealing treatment can decrease aliquation and agglomeration of the dendritic eutectic structure. 3.1.2 Al3Ni phase analysis Using metallographic analysis methods, we calculated the Al3Ni plates’ average size class of material A, and B in
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Gao Peng et al. / Rare Metal Materials and Engineering, 2013, 42(1):0006-0013
Table 4 a
Average size of Al3Ni phase in different annealing time (μm)
b Material
0h
24 h 48 h 75 h
100 h
150 h
200 h
A
11.5
10.6
8.6
5.8
3.9
3.2
3.1
B
18.5 16.9 13.3
11.6
9.8
9.2
8.4
b a c
d
Al3Ni phase
25 µm
Fig.5
OM images of material A (a, b), and B (c, d) with annealing treatment: (a,c) 24 h, and (b, d) 200 h
2σ Vm r
μ = G m + (1 − x )
∂G m ∂x
Fig.6
Simplified model of Al3Ni phase before fusing and spheroidizing
different annealing time (for every different status we chose more than ten photos to statistic, and the plates’ statistic result showed the average length of major axis and minor axis), shown in Table 4. The average size of Al3Ni plates became smaller as the annealing time extended. When the annealing time extended to 200 h, the average size of Al3Ni plates can reach to 3.1 μm in material A and 8.4 μm in material B. This could illustrate long time annealing treatment can decrease the average size of Al3Ni plates clearly. And during 24 to 100 h, the effects of refining Al3Ni plates are more significant. This could be found both in the specimens with annealing and with annealing + aging. 3.1.3 Spheroidizing mechanism analysis In order to explain the reasons for the large Al3Ni plates fusing and spheroidizing to small plates after annealing, we built the model in Fig.6. Fig.6 is the simplified diagram of Al3Ni plates before fusing and spheroidizing. The curvature radius of point a, b is ra and rb. Obviously, ra is much less than rb. By the thermodynamic formula [8] G m = G m (0) +
Matrix
(1) (2)
Gm (0): The molar Gibbs free energy when the curvature radius is zero r: The curvature radius σ: The surface tension Vm: The molar volume X: The mole fraction of group element The chemical potential difference of two points was:
∂Gm(a) ∂Gm (b) − (1 − Xb ) ∂Xa ∂Xb (3) ∂Gm (a) ∂Gm (b) 2σ Vm 2σ Vm = − + (1 − X a ) − (1 − X b ) ra ∂X a ∂X b rb
Δμ = Gm(a) − Gm (b) + (1 − Xa )
Because Xa = Xb
Δμ =
2σ V m 2σ V m − ra rb
(4)
Since ra is much less than rb, the chemical potential at point a is higher than that at point b in the matrix. According to Ref.[9], at 400 ºC, the Ni atoms have been able to spread long distance, so it is enough for Ni atoms to carry out a long-ranged diffusion at 470 ºC. The Ni atoms would diffuse from point a with higher chemical potential to the point b with lower chemical potential[10], which brought damage to the equilibrium concentration of Ni atoms near the point a. In order to keep balance, the Al3Ni dissolved continuously to make up the shortage of nickel concentration, so the groove would be deepening and widening continuously, the sharp corner would be dissolving constantly, until the branch had fused, and the sharp corner had disappeared and spheroidized. The fusing of the branch is bound to create new grooves and sharp corners, and then it would continue this process. After repeated fusing of branches and spheroidizing of sharp corners, we eventually obtained tiny Al3Ni grains which were nearly spheroidal, see in Fig.5b.
3.1.4
Hardness test
We tested the hardness after annealing and T6 aging treatment (the load was 1 kg, and load-on 30 s. Average value from ten specimens), the results were shown in Table 5 and Fig.7.
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Gao Peng et al. / Rare Metal Materials and Engineering, 2013, 42(1):0006-0013
Table 5
Results of Vickers hardness test (HV) after annealing(MPa)
Annealing time
Sample A
Sample B
7050
0 h+T6
1918
2364
180
24 h+T6
1980
2385
—
48 h+T6
2050
2385
—
75 h+T6
2080
2395
—
100 h+T6
1957
2293
—
150 h+T6
1952
2278
—
200 h+T6
1950
2272
—
Hardness, HV/MPa
2400
Material A Material B
2300 2200 2100 2000 1900
0
50
100
150
200
Annealing Time/h Fig.7
Curves of hardness-annealing time
phase is the main reason of fracture. So after annealing treatment experiments, we tried to use traditional deformation method to lower the size of the large pro-eutectic phase and reduce its brittleness. In following experiments, Cold-Rolling and Hot-Rolling were adopted to smash the Al3Ni plates. The original thickness of sample was 10 mm, every time the rolling reduction is 1 mm, and after seven-step rolling, the total deformation rate was 70%. The microstructure of Al3Ni phases in different deformation is shown in Fig.8. (H-R stands for Hot-Rolling and C-R stands for Cold-Rolling in the following.) The microstructure showed that, when the C-R deformation rate was 10%, a few large Al3Ni plates were fractured. Al3Ni long plates which were parallel to the rolling stress need large deformation and the possibility of fracture is big, while the long plates which were perpendicular to the rolling stress can have compressive deformation to offset the deformation. Most of the Al3Ni plates were longer than 50 μm, the average space between Al3Ni plates was large, about 40 μm, and some Al3Ni plates had cracks in their cusps. The H-R showed similar situation to C-R, but most of the Al3Ni plates were integrated, cracks can be easily found but clear fractures had not existed. Because at the high temperature (450 ºC), Al-matrix was ductile and easy to deform, the Al-matrix had compressive deformation which was the main effect of the compressive stress, and Al3Ni plates mainly had position change.
The hardness of the aluminum matrix composite alloy were much higher than that of 7050 alloy, and increased with the content of Ni, which indicated that the hardness of 7050 alloy could be significantly improved by Al3Ni phases. From the Fig.7, with the prolonging of annealing time, the hardness first increased and then decreased, especially for the material A who has near eutectic composition. This phenomenon may be caused by the following reasons: (1) the casting defects of the materials (such as porosity, segregation, dispersed shrinkage, shrinkage cavity and so on) were eliminated by the annealing treatment, which made the structure more compact; (2) the Al3Ni is hard and brittle phase, so their presence might significantly improve the hardness of materials, especially when they were dispersed in Al matrix, and it is particularly reflected on the material A, whose hardness was improved much more than that of the material B; (3) as the annealing time was further extended, the dispersion effect of Al3Ni phase on the increase of hardness would not be strengthened, but the matrix would have the softening trend because of the elimination of residual stress, which led to the decline in hardness. 3.2 Rolling treatment and research 3.2.1
Cold-Rolling/Hot-Rolling
From the former part, we know the break of large Al3Ni
a
b
c
d
e
f
50 µm
Fig.8
Microstructure of Al3Ni after H-R/C-R: (a) 10%C-R, (b) 40%C-R, (c) 70%C-R, (d) 10%H-R, (e) 40%H-R, and (f) 70%H-R)
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Gao Peng et al. / Rare Metal Materials and Engineering, 2013, 42(1):0006-0013
Average Length/µm
11
a
10
Minor axis of C-R Minor axis of H-R
9 8 7 6 5 4 3 10
20
30
40
50
60
120
70 b
Major axis of C-R Major axis of H-R
100 Average Length/µm
During deformation rate change from 20% to 40%, the Cold-Rolling and Hot-Rolling showed the same trend, the Al3Ni plates became longer continuously. When the deformation was about 40%, the Al3Ni plates reach the extreme of length direction extension, and long Al3Ni plates were inclined to rupture owing to the shearing stress. When the deformation rate is 50%, long Al3Ni plates were smashed into short plate of which the major axis to minor axis ratio is about 5:1; at that time, most of the long Al3Ni plates had disappeared, breakage effect of the compressive stress began to diminish but the compressive effect of Al3Ni short plate began to increase. During the deformation rate change from 60% to 70%, the short Al3Ni plates attenuated and became secondary long Al3Ni plate, then obvious rupture appeared again, and finally the Al3Ni plates were crushed into grains of which the major axis to minor axis ratio is about 2:1. Different from cold-rolling, Al3Ni grains of 70% deformation after hot-rolling had clearer cusps than that after cold-rolling. It is because during the hot rolling, the yield strength of Al-matrix is low and Al-matrix was ductile, which transmit less stress to Al3Ni phase than that under cold-rolling, so the crushing effect was less obvious. Under the cold-rolling condition, Al3Ni plates can squeeze each other to eliminate the cusps, but under the hot-rolling condition, the stress between Al3Ni plates could be transmitted to Al-matrix and cause Al-matrix deformation, so cusps were preserved finally. 3.2.2 Al3Ni phase analysis Since the Al3Ni phase shape is mainly plate, we made the hypothesis that the cross-section of Al3Ni phase can be considered as symmetrical rectangle, four photos of each rolling state were chosen for data analysis and 128 Al3Ni plates were measured in the four photos, measuring objects included the lengths of major axis and minor axis of Al3Ni plates. And the curve of the Al3Ni plates’ average length in different deformation rates was made and is shown in Fig.9. The curves in Fig.9a show that both the H-R and C-R can reduce the length of minor axis. After 70% deformation of H-R, the average minor axis is about 4.4 μm comparing to 3.4 μm of C-R. The curves also indicated that, with the increase of deformation rate under C-R, the length of minor axis of Al3Ni phase decreased continuously, but when the deformation rate is 60% under H-R, the minor axis did not attenuate obviously after 7th hot rolling and maintained 4.4 μm. during the 7th hot rolling, the main effect of compressive stress was to smash the Al-Ni plate further, and the average length of major axis was reduced from 37.5 to 25.8 μm Since the influence of rolling on the minor axis of Al3Ni phase is inclined to be linear, curves were fitted as below(X represents deformation rate and Y represents the length of minor axis).
80 60 40 20 0 10
20
30
40
50
60
70
Deformation Rate/% Fig.9
Curves of the Al3Ni plates’ average length in different deformation rates of cold-rolling (C-R) and hot-rolling (H-R)
YC-R=–11.89X+ 11.51 YH-R=–7.607X + 9.214 It can be found that the reducing effect of minor axis length of cold-rolling is 56% better than that of hot-rolling. Fig.9b indicates that there exist peak value of the major axis length during the multi-step hot rolling process. When the deformation rate is 30%, the average length reached the maximum of 106 μm. From Fig.8, we know the Al 3Ni plates surfaces were full of cracks but not completely fractured, but when the deformation rate reached 40%, cracks were enlarged and one long plate became several short plates, and the average length of major axis decreased to 80 μm. From the OM photos, fractures mainly took place on long Al3Ni plates; short Al3Ni plates attenuated but did not fracture. Different from the H-R, there were two peak values of major axis length during the multi-step C-R process; they were 105 μm of 30% deformation and 37.6 μm of 60% deformation. After 30% deformation, there were many cracks which inclined to the major axis at an angle of 45º but did not fracture completely. After 40% deformation, the average major axis length suddenly dropped to 40 μm, most the long Al3Ni plates had been smashed into secondary Al3Ni plate, and the major axis to minor axis ratio is about
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6.4:1, close to that of the 10% deformation condition. When the deformation rate is 60% which, another peak value of major axis length appeared, was 37.6 μm. It is because the focus of compressive stress in Al3Ni plates is the determinants. 3.2.3 Hardness test We tested the hardness after every step of cold/hot-rolling treatment (the load was 1 kg, and load-on 30 s), and each specimen tested ten points and calculated the average value. The results were shown in Table 6. Then we studied the hardness curve of Al-matrix (average value from ten specimens) in Cold-Rolling and Hot-Rolling process, with microhardness instrument to avoid the Al3Ni phases (the load was 0.0 1kg, load-on 10 s). showed in Table 7. As the deformation rate increased, the hardness of material B increased obviously both in Cold-Rolling or Hot-Rolling process. And at 30%, 60% deformation rate, their hardening effect is not such obviously. In Cold-Rolling process, work-hardening effect is the main reason of hardness increased. And in Hot-Rolling process, the hardness increased and with a peak value. After the 30% deformation rate, the soften effect was larger. So the hardness of Al-matrix had fallen down. The hardness of material B increased with deformation rate. But at 60% (C-R) and 30% (H-R) deformation rate, there was a minimum value, and it was corresponding with the change of Al3Ni plates’ size. In the former condition, the Al3Ni plates became larger, so the precipitation hardening effect was significantly weakened by the size effect. Table 6
Results of Al-Ni phases’ Vickers hardness test (HV) after C-R/H-R(MPa) 10%
20%
30%
40%
50%
60%
70%
C-R
921
969
974
992
1041
1015
1090
H-R
925
932
947
935
937
955
1032
Table 7
Results of Vickers hardness (HV) in Al matrix after multi-step C-R/H-R(MPa) 10%
4
20%
30%
40%
50%
60%
70%
C-R
686
705
708
714
757
776
807
H-R
682
708
740
722
695
674
661
Discussion and Conclusions
1) From Ref.[13,14], the effect on mechanical property of the material after Al3Ni phase generated could be illustrated as the following formula: First, the effect of Al3Ni plates on strength can be expressed as σ y = σ mVm +σ r Vr r
(5)
σy: The calculated valued of yield strength σm, σr: The yield strength of matrix and reinforcement particle Vm, Vr: The volume fraction of matrix and reinforcement particle And actually, the strength of material is not only affected by the volume fraction. When the size of reinforcement particle is large enough, the effect of Orowan bypassing dislocation strengthening mechanism would be ignored. In the external stress, the slipping of dislocations in matrix was slowed down at interface of matrix/Al3Ni plates, with stress concentration in the Al3Ni plates. The stress value σy could be described as [14]:
σy=
3G m G p b V p 2 d (1−V p )c
(6)
Gm, Gp: the shear modulus of matrix and Al3Ni plate b: Burgers vector Vp: the volume fraction of Al3Ni plate c: constant d: average size of Al3Ni plates By substituted relevant parameters, we can infer that when the average size of Al3Ni phase reaches to 5 μm and the volume fraction of Ni reaches to 30%Al3Ni, the yield strength of the material can get to 630 MPa, far more than 500MPa of 7050 aluminum alloy. 2) Through the fracture analysis, the fracture mechanism of the material is mainly Al3Ni fracture and interface debonding. And the plasticity decreased as the content of Ni increased. 3) As the annealing time extending, the Al3Ni phase had a trend of fusing and spheroidizing. After long time annealing, the aluminum matrix composites could obtain tiny and dispersed Al3Ni grains which were nearly spheroidal. The average diameter of the Al3Ni grains is about 3 μm in the material A (5%Ni), and it is 6 μm in the material B (10%Ni). 4) Besides annealing, multi-step Hot-rolling/Cold-rolling process also can make the Al3Ni plates size decrease obviously. The minor axis length of Al3Ni plates decreased continuously due to the compressive effect and smashing effect, and the length in minor axis at 70% deformation rate is about 4 μm. Comparing to the minor axis, there existed peak values in major axis length, mainly because of the compressive effect which squeeze the Al3Ni plates to be longer before the stress reached to the threshold value. When the stress reached to the threshold value, the Al3Ni plates fractured and then the major axis length dropped rapidly. From statistical results, the effect of Cold-rolling was slightly better than that of hot-rolling. 5) The addition of Ni can significantly improve the hardness of 7050 aluminum alloy. After T6 aging, the
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Gao Peng et al. / Rare Metal Materials and Engineering, 2013, 42(1):0006-0013
hardness of material A (5%Ni) and material B (10%Ni) were 1918 and 2364 MPa respectively, and both were much higher than 180 HV, the hardness of the 7050 deformed aluminum alloy. As the annealing time extending, the hardness of all first increased and then decreased, especially for material A (5%Ni), whose hardness could be up to 2070 MPa after 48 h annealing and T6 aging, and the hardness of material B (10%Ni) could be up to 2385 MPa after 48 h annealing and T6 aging.
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