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Mesoscopic structure control of spray formed high strength Al–Zn–Mg–Cu alloys
Acta Materialia 53 (2005) 2919–2924 www.actamat-journals.com
Mesoscopic structure control of spray formed high strength Al–Zn–Mg–Cu alloys M.M. Sharma
a,*
, M.F. Amateau b, T.J. Eden
c
a
b
Bucknell University, Department of Mechanical Engineering, Dana Building, PO Box A0551, Lewisburg, PA 17837, USA The Pennsylvania State University, Department of Engineering Science and Mechanics, 212 EES Building, University Park, PA 16802, USA c Applied Research Laboratory, The Pennsylvania State University, PO Box 30, N. Atherton St, State College, PA 16804-0030, USA Received 29 May 2004; received in revised form 10 February 2005; accepted 4 March 2005 Available online 21 April 2005
Abstract Several high solute, high strength 7xxx series aluminum alloys with solute contents close to equilibrium solid solubility limits of the Al–Zn–Mg–Cu system have been produced by rapid solidification using spray deposition (the Osprey process). This process yields massive preforms directly from the liquid state by combining atomization and consolidation into one step. Various elements, including chromium, manganese and silver are incorporated to produce a variety of microstructures and mechanical properties. The zinc to magnesium ratio is also varied to see the effect on the strength. Superior strengths in excess of 849 MPa are achieved and are attributed to two major substructures with different scale; nanometer sized g 0 metastable precipitates and slightly larger, but finely distributed dispersoids which provide a fiber-like reinforcement. The remarkable strengthening is predominantly attributed to precipitation hardening and a large coherency strain. 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Atomization; Aluminum alloys; Precipitation; Small angle neutron scattering; Transmission electron microscopy
1. Introduction Several attempts have been made to improve the mechanical properties of Al–Zn–Mg–Cu alloys [1,2]. Much of the progress achieved in the development of aluminum alloys can be attributed to new production methods, such as powder metallurgy (PM) and spray deposition. These methods utilize the rapid solidification process, which allows increased compositional latitude, while avoiding common problems associated with conventional ingot metallurgy, such as alloy segregation and microstructural coarsening. The success of spray deposition is related to faster cooling rates and thus preferred microstructures that produce superior mechanical properties. Finer microstructures and reduced microseg*
Corresponding author. Tel.: +1 570 577 1686; fax: +1 570 577 7281. E-mail address: [email protected] (M.M. Sharma).
regation can lead to reduced homogenization times during solution heat treatment and a variety of mesoscopic structure changes which lead to superior mechanical properties. The successful exploitation of this technology has prompted the development of a diverse group of aluminum alloys with unparalled combinations of strength and thermal stability. The mechanical properties of Al–Zn–Mg–Cu alloys are defined in terms of precipitation hardening, where nanometer sized Guinier–Preston zones and/or metastable g 0 precipitates act as pinning centers for dislocation movement. Improvement in mechanical properties will be achieved by controlling the size, shape, distribution and spacing of precipitates. A systematic investigation was carried out using various types of Al–Zn–Mg–Cu based alloys. The control of the mesoscopic structure using various chemical compositions and reasons for achieving such high strength is discussed theoretically.
1359-6454/$30.00 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2005.03.007
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M.M. Sharma et al. / Acta Materialia 53 (2005) 2919–2924
2. Experimental procedure
Table 1 Alloy designations and compositions (in wt.%)
2.1. Material processing
Alloy designation
Zn
Mg
Cu
Mn
9Zn–0.0Cu–0.36Mn–0.28Cr 9Zn–1.3Cu–0.34Mn–0.22Cr 12Zn–1.1Cu–0.29Mn–0.02Cr 12Zn–1.1Cu–0.18Mn–0.15Cr 12Zn–1.1Cu–0.20Mn–0.16Cr 8Zn–1.6Cu–3.97Mn–0.04Ag
9.77 9.31 12.00 12.00 12.00 8.21
3.03 3.49 3.28 3.25 3.34 2.78
0.01 1.25 1.15 1.14 1.12 1.58
0.36 0.34 0.29 0.18 0.20 3.97
The main features of the Osprey process are described elsewhere [3]. The spray deposition method differs from that of PM technology by combining the atomization and consolidation steps into one operation. As a result, it is possible to produce very dense, almost net-shaped preforms of rapidly solidified materials, directly from the liquid state as shown in Fig. 1. In this study, binary master alloys and elemental metals were obtained from several commercial sources. The spray deposition experiments were conducted in a specially designed spray chamber. First, 30–70 pounds of alloy was superheated to temperatures between 850 and 1000 C in a fiber crucible using an induction apparatus. Second, the superheated alloy melt was delivered to an atomizer where it was atomized into a fine dispersion of 5–40 lm-sized droplets using nitrogen gas. After atomization, the distribution of partially solidified droplets was collected on a rotating substrate (rotation 3 Hz) positioned at approximately 700 mm from the gas atomizer (scan speed 19–21 Hz). To maintain the atomizer-substrate distance, the substrate was displaced vertically during spraying at a withdrawal rate of 0.7–1.04 mm/s. In order to obtain billets of different diameters, the scanning mechanism was adjusted to increase or decrease the outer scan radius. The selection of the atomizer–substrate distance
and other processing variables used was prompted by experimental and numerical results presented elsewhere [4]. These results indicated that at this flight distance, the distribution of atomized droplets contained liquid, semisolid and solid droplets [5]. 2.2. Thermal mechanical processing and heat treatment Six alloys of various compositions were processed by spray metal forming for this investigation. Table 1 displays the chemical composition and designation of the alloys. Spray formed billets of each alloy composition were then hot extruded at 400 C into round rod at an extrusion ratio of 44:1. Alloys were then solution treated between 470 and 490 C for 0.5–2 h, cold water quenched and then peak aged to a T6 heat treatment. The extruded and heat treated rods were then cut into bars 76.2 mm (3.0000 ) in length to make sub-size tensile specimen. The bars were then turned down on a lathe to a nominal diameter of 6.35 mm (0.25000 ) in the test section. The grip diameter was 7.62 mm (0.30000 ) with threaded ends to prevent slipping while testing. A minimum fillet radius of 3/1600 was used, with a reduced section at least 31.75 mm (1.2500 ) in length. The samples were machined according to specifications stated in ASTM E8-95a, for sub-size specimens. A minimum of five samples per alloy composition were used for uniaxial tensile testing. The microstructure of the extruded and heat treated alloys was observed by means of transmission electron microscopy and small-angle neutron scattering (SANS).
3. Results and discussion
Fig. 1. A schematic cross-section of the spray metal forming process showing how atomization and consolidation are combined into one step.
When a glide dislocation encounters an array of precipitates, it bows out with a curvature, at some critical angle 0 6 uc P p, between two precipitates before it can move on. The angle uc is a measure of the strength of the obstacle. According to Brown and HamÕs analysis [6], the critical shear stress at which a dislocation can move a large distance through an array of precipitates is a function of precipitate spacing (L) on the slip plane and the critical angle (uc). Weak precipitates can be overcome with very slight bending (uc = 180), whereas
M.M. Sharma et al. / Acta Materialia 53 (2005) 2919–2924
a strong precipitate cannot be overcome unless the dislocation practically doubles back on itself (uc = 0). In the latter case, the yield stress is defined by
1=2
;
ð2Þ
where G is the shear modulus, e is the coherency strain, l is the average radius and Vf is the volume fraction of precipitates. Given this information, the coherency strain for the present spray formed alloys with respect to volume fraction of metastable precipitates was estimated. First, in order to find the strengthening attributed only to the precipitation of structural particles, the change in yield strength is needed, and is defined as, Dry ¼ robs rAsQ ;
Alloy designation
Rp (nm)
Vf
(Rp * Vf)1/2
Yield strength (MPa)
9Zn–0.0Cu–0.36Mn–0.28Cr 9Zn–1.3Cu–0.34Mn–0.22Cr 12Zn–1.1Cu–0.29Mn–0.02Cr 12Zn–1.1Cu–0.18Mn–0.15Cr 12Zn–1.1Cu–0.20Mn–0.16Cr 8Zn–1.6Cu–3.97Mn–0.04Ag
2.42 2.53 2.61 2.71 2.73 2.73
0.047 0.049 0.059 0.064 0.064 0.065
0.337 0.352 0.392 0.416 0.418 0.421
711 794 807 800 842 849
ð1Þ
and corresponds to the Orowan stress, where T is the line tension, M is the Taylor factor and b is the Burgers vector. Under a simple elastic model, the line tension is given as T = Gb2/2, where G is the shear modulus of the precipitate and M is 3.06 for the polycrystalline facecentered cubic metals. Conversely, the weak interaction between precipitate and dislocation is called a cutthrough mechanism. Several models exist which describe this behavior. The coherency strain effect has been proven to be dominant for G.P. zones and metastable precipitates in the Al–Zn based alloys [2,7–12]. As a result, the contribution to strengthening from misfitting metastable precipitates was estimated in this study. The strengthening of alloys by misfitting coherent particles occurs by virtue of the interaction between stress fields of the precipitates and the dislocation. Due to the fact that the dislocation interacts the strain field of the precipitates that physically intersect its slip plane as well as the strain fields of precipitates that do not, it is necessary to find some means of averaging the effects of all the particles in the alloy. The methods of averaging the effects of all the precipitates in the microstructure on the yield strength have been discussed by Brown and Ham [6] and Gerold and Haberkorn [8], with only the constant differing from one theory to the other. According to Gerold and Haberkorn [8] and other researchers [9,10], the yield stress is expressed by the following equation: rys ¼ 3MGe3=2 ðlV f =bÞ
Table 2 Measured structure parameters for the spray formed alloys
ð3Þ
where the as-quenched yield strength is subtracted from the observed yield strength. The change in yield strength is thus the increment due to the formation of structural precipitates only. This value can then be plotted versus the structure parameters to evaluate the relationship between strength and structure of the materials. A summary of the measured structure parameters with the corresponding yield strength value is presented in Table 2. RP is the precipitate radius and the volume fraction of the precipitates is Vf. Using this information, the yield strength of these materials can be explained as a
function of the average radius and volume fraction of the structural precipitates. As-quenched data was obtained in this study, but was not considered reliable, due to the extremely fast precipitation rate during quenching, especially in the alloys with silver addition. However, if the as-quenched data is assumed to be approximately the same for all specimens, it is possible to estimate the coherency strain. Using Eq. (2), while all values other than precipitate radius (RP) and volume fraction (Vf) remain constant, the yield strength versus the square-root of the precipitate radius and volume fraction of precipitate was plotted, Fig. 2. The slope of the line is the fitted result of Eq. (2). The coherency strain is thus estimated using G = 26.2 GPa [13], b = 0.29 nm [6], M = 3.06 [7] and the slope of the line. Furthermore, the intercept from the equation of the slope of the line can be used as an approximation for the as-quenched yield strength. Table 3 presents the experimental as-quenched data for the spray formed alloys. Higher values obtained for the asquenched yield strength compared to the calculated value is attributed to the fast precipitation rate in these alloys. The time between solution treatment temperature and quench, even though very short, can allow for the beginning of precipitation of structural particles. This can make 900
y = 1244.3x + 319.94 850 Yield Strength (MPa)
rOrowan ¼ 1.6MT =bL
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800 750 700 650 600 0.250
0.300
0.350 Rp(Vf)
0.400 1/2
0.450
0.500
1/2
(nm )
Fig. 2. Change in the observed yield strength as a function of the average radius and volume fraction of g 0 precipitates. The slope can be described by the term of precipitation hardening.
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M.M. Sharma et al. / Acta Materialia 53 (2005) 2919–2924
Table 3 As-quenched yield strength data for the spray formed alloys Alloy designation
As-quenched yield strength (MPa)
9Zn–0.0Cu–0.36Mn–0.28Cr 9Zn–1.3Cu–0.34Mn–0.22Cr 12Zn–1.1Cu–0.29Mn–0.02Cr 12Zn–1.1Cu–0.18Mn–0.15Cr 12Zn–1.1Cu–0.20Mn–0.16Cr 8Zn–1.6Cu–3.97Mn–0.04Ag
341 350 320 335 349 379
the experimental value for as-quenched strength slightly higher than the actual value. Table 4 presents the value for estimated coherency strain of the Al–Zn–Mg–Cu spray formed alloys and for other Al–Zn based alloy systems. From this estimation, it appears that the Al–Zn–Mg–Cu alloy system has a greater coherency strain than other systems, and processing these alloys through spray forming results in a further increase in coherency strain. The estimated value for coherency strain compares well with that of P/M processed high solute alloys. However, the data from that study was for high manganese alloys, alone. As a result, a slightly higher coherency strain was realized in those alloys. Nevertheless, the estimation validates the theory of contribution to strengthening by misfitting precipitates. Gerold and Haberkorn [8] established that after the precipitate exceeds a certain size r, the increment in flow-stress, s, with further aging increases with volume fraction of precipitate in proportion with (Vf)1/2. They explained this behavior with the help of data from Dash and Fine [11]. It was considered that Vf was essentially constant after a specific time of aging. However, the flow-stress continues to increase with aging time, despite the fact that Vf remains constant. During aging, as the particle coarsens, the average particle radius becomes larger obeying the power law r = f(t)1/3, while the volume fraction remains constant. Then, the structure parameters can be explained in terms of the Orowan stress expressed in Eq. (1), rewritten as rOrowan ¼ ð0.553MGbðV f Þ1=2 Þ=r;
ð4Þ
where r is the precipitate radius. In this equation, the Orowan stress [6,14] will decrease with increasing precipitate radius, while the yield stress from the cutthrough model will continue to increase (Eq. (2)). Using these equations, Osamura et al. [12] proposed that the theoretical maximum yield strength of these alloys would be found to be at the cross-over point for both
Fig. 3. Bright field (a) and dark field (b) electron photomicrograph of massive structural precipitation in the 12Zn–1.1Cu–0.20Mn–0.16Cr spray formed alloys.
these mechanisms. It was also proposed that the yield strength would increase with increasing coherency strain and volume fraction of precipitate, and that the strength of 1 GPa could theoretically be achieved in high solute
Table 4 Coherency strain (e) for several Al–Zn based alloy systems Al–Zn
Al–Zn–Mg
Al–Zn–Mg–Cu
Al–Zn–Mg–Cu
Al–Zn–Mg–Cu
[8] 0.007
[2] 0.015
Commercial [10] 0.022
P/M-high solute [8] 0.031
Spray formed-high solute 0.029
M.M. Sharma et al. / Acta Materialia 53 (2005) 2919–2924
concentration alloys. The solution for the theoretical yield stress at this cross-over point according to Osamura et al. [12] is as follows: rys;max ¼ 1.71MGeðV f Þ1=2 :
ð5Þ
It is clear from this analysis that three structure parameters, coherency strain, volume fraction and shear modulus, are important to the strengthening of high solute alloys. In this study, the spray formed alloys contained a variety of solute elements and an increased amount of zinc content compared to conventional 7xxx series alloys. As a result, a high volume fraction of metastable precipitates resulted in the spray formed materials and are shown in Fig. 3. The variations of amounts of zinc, magnesium, copper and silver content contributed to resulting microstructure. The estimation of the high coherency strain achieved in the spray formed alloys can be explained as follows. Chromium or manganese additions promote the formation of the rod-like E-phase or the quasi-ternary (Al20(Cu,Zn)2Mn3) intermediate phases, respectively, thus providing a fiber-like reinforcement Fig. 4. The zinc is distributed in both the E- and quasi-ternary phases, in addition to the g 0 phase. However, the magnesium is distributed in great amount in only the g 0 phase, while the E-phase has some distribution of magnesium and the quasi-ternary phase contains none. The high magnesium content tends to generate a high coherency strain due to the atomic size difference with aluminum. The microstructural characteristics of the alloys can be used to rationalize the observed mechanical properties. The estimation of coherency strain was performed in this work to correlate the increase in yield strength
Fig. 4. Bright-field photomicrographs of the fiber-like dispersoid Ephase in the 9Zn–0.0Cu–0.36Mn–0.28Cr alloy.
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Table 5 Measured structural particle concentration for various compositions Alloy designation
Particle concentration
9Zn–0.0Cu–0.36Mn–0.28Cr 9Zn–1.3Cu–0.34Mn–0.22Cr 12Zn–1.1Cu–0.29Mn–0.02Cr 12Zn–1.1Cu–0.18Mn–0.15Cr 12Zn–1.1Cu–0.20Mn–0.16Cr 8Zn–1.6Cu–3.97Mn–0.04Ag
3.0 · 104 3.5 · 104 4.0 · 104 4.0 · 104 4.0 · 104 4.5 · 104
T6
with the dependence of the size and volume fraction of metastable precipitates. This is attributed to the increasing zinc content of the material. For the most part, with similar processing conditions, increasing the zinc content resulted in an increased volume fraction of metastable precipitates, Table 5. Considering the high coherency strain estimated for the high solute spray formed alloys, and the measured high volume fraction of metastable precipitates, it can be hypothesized that the increased addition of zinc, and thus a high volume fraction of structural particles, results in an increased coherency strain. However, to verify this theory, the actual coherency strain must be measured. Measuring the coherency strain of fine precipitates cannot effectively be measured using transmission electron microscopy, since the misfit strain is often below 1% and the error is approximately 5%.
4. Summary The characteristic microstructure of the spray formed alloys is predominantly associated with a massive precipitation of g 0 plates. This microstructure is most easily achieved with alloys containing a high amount of solute content, with Zn + Mg contents over 15 wt.% where the zinc content is at least 12 wt.%. A massive and fine distribution of the g 0 phase can also be achieved with lower amounts of Zn and Mg, if accompanied by the addition of manganese and minor amounts of silver (less than 0.04 wt.%). In this study, the addition of manganese precipitated the rod-like quasi-ternary phase, which provided additional strengthening through fiber reinforcement. Therefore, it was concluded the ultrahigh strengths of the spray formed alloys is related to multiple effects of both precipitation hardening and fiber reinforcement. The reason for increased yield strength contributed from only precipitation hardening in the spray formed alloys can be explained in terms of coherency strain and a coherency strain model. The coherency strain was estimated using a model and found to be much higher in the spray formed alloys compared to PM or commercial alloys of similar composition.
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References [1] Pickens JR. J Mater Sci 1981;16:1437–57. [2] Osamura K et al. Development of high strength aluminum alloys by mesoscopic structure control. Metall Trans A 1995;26:1597–9. [3] Leatham AG, Brooks RG, Yamen M. Modern developments in powder metallurgy. Princeton, NJ: MPIF Press; 1985. pp. 15–17 and 157–169. [4] Bewlay BP, Cantor B. J Mater Res 1991;6:1433. [5] Grant PS, Kim WT, Cantor B. Mater Sci Eng A 1991;134: 1111.
[6] Brown LM, Ham RK. In: Kelly A, Nicholson RB, editors. Strengthening methods in crystals. Amsterdam: Elsevier; 1971. p. 12. [7] Mott NF, Nabarro FRN. Proc Phys Soc 1940;52:86. [8] Gerold V, Harberkorn H. Phys Stat Sol 1966;16:675–9. [9] Osamura K, Okuda H, Ameniya Y, Hashizume H. Metall Trans A 1988;19:1973–8. [10] Osamura K, Ochiai S, Uehara T. Jap J Light Metal 1984;34:192. [11] Dash, Fine. Mater Sci Form 1993;342–343:1345–67. [12] Osamura K et al. Mater Sci Form 1966;217–222:1829–34. [13] Wilm A. Metallurgie 1911;8:225. [14] Orowan E. Z Phys 1934;89:634.
Mesoscopic structure control of spray formed high strength Al–Zn–Mg–Cu alloys M.M. Sharma
a,*
, M.F. Amateau b, T.J. Eden
c
a
b
Bucknell University, Department of Mechanical Engineering, Dana Building, PO Box A0551, Lewisburg, PA 17837, USA The Pennsylvania State University, Department of Engineering Science and Mechanics, 212 EES Building, University Park, PA 16802, USA c Applied Research Laboratory, The Pennsylvania State University, PO Box 30, N. Atherton St, State College, PA 16804-0030, USA Received 29 May 2004; received in revised form 10 February 2005; accepted 4 March 2005 Available online 21 April 2005
Abstract Several high solute, high strength 7xxx series aluminum alloys with solute contents close to equilibrium solid solubility limits of the Al–Zn–Mg–Cu system have been produced by rapid solidification using spray deposition (the Osprey process). This process yields massive preforms directly from the liquid state by combining atomization and consolidation into one step. Various elements, including chromium, manganese and silver are incorporated to produce a variety of microstructures and mechanical properties. The zinc to magnesium ratio is also varied to see the effect on the strength. Superior strengths in excess of 849 MPa are achieved and are attributed to two major substructures with different scale; nanometer sized g 0 metastable precipitates and slightly larger, but finely distributed dispersoids which provide a fiber-like reinforcement. The remarkable strengthening is predominantly attributed to precipitation hardening and a large coherency strain. 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Atomization; Aluminum alloys; Precipitation; Small angle neutron scattering; Transmission electron microscopy
1. Introduction Several attempts have been made to improve the mechanical properties of Al–Zn–Mg–Cu alloys [1,2]. Much of the progress achieved in the development of aluminum alloys can be attributed to new production methods, such as powder metallurgy (PM) and spray deposition. These methods utilize the rapid solidification process, which allows increased compositional latitude, while avoiding common problems associated with conventional ingot metallurgy, such as alloy segregation and microstructural coarsening. The success of spray deposition is related to faster cooling rates and thus preferred microstructures that produce superior mechanical properties. Finer microstructures and reduced microseg*
Corresponding author. Tel.: +1 570 577 1686; fax: +1 570 577 7281. E-mail address: [email protected] (M.M. Sharma).
regation can lead to reduced homogenization times during solution heat treatment and a variety of mesoscopic structure changes which lead to superior mechanical properties. The successful exploitation of this technology has prompted the development of a diverse group of aluminum alloys with unparalled combinations of strength and thermal stability. The mechanical properties of Al–Zn–Mg–Cu alloys are defined in terms of precipitation hardening, where nanometer sized Guinier–Preston zones and/or metastable g 0 precipitates act as pinning centers for dislocation movement. Improvement in mechanical properties will be achieved by controlling the size, shape, distribution and spacing of precipitates. A systematic investigation was carried out using various types of Al–Zn–Mg–Cu based alloys. The control of the mesoscopic structure using various chemical compositions and reasons for achieving such high strength is discussed theoretically.
1359-6454/$30.00 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2005.03.007
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M.M. Sharma et al. / Acta Materialia 53 (2005) 2919–2924
2. Experimental procedure
Table 1 Alloy designations and compositions (in wt.%)
2.1. Material processing
Alloy designation
Zn
Mg
Cu
Mn
9Zn–0.0Cu–0.36Mn–0.28Cr 9Zn–1.3Cu–0.34Mn–0.22Cr 12Zn–1.1Cu–0.29Mn–0.02Cr 12Zn–1.1Cu–0.18Mn–0.15Cr 12Zn–1.1Cu–0.20Mn–0.16Cr 8Zn–1.6Cu–3.97Mn–0.04Ag
9.77 9.31 12.00 12.00 12.00 8.21
3.03 3.49 3.28 3.25 3.34 2.78
0.01 1.25 1.15 1.14 1.12 1.58
0.36 0.34 0.29 0.18 0.20 3.97
The main features of the Osprey process are described elsewhere [3]. The spray deposition method differs from that of PM technology by combining the atomization and consolidation steps into one operation. As a result, it is possible to produce very dense, almost net-shaped preforms of rapidly solidified materials, directly from the liquid state as shown in Fig. 1. In this study, binary master alloys and elemental metals were obtained from several commercial sources. The spray deposition experiments were conducted in a specially designed spray chamber. First, 30–70 pounds of alloy was superheated to temperatures between 850 and 1000 C in a fiber crucible using an induction apparatus. Second, the superheated alloy melt was delivered to an atomizer where it was atomized into a fine dispersion of 5–40 lm-sized droplets using nitrogen gas. After atomization, the distribution of partially solidified droplets was collected on a rotating substrate (rotation 3 Hz) positioned at approximately 700 mm from the gas atomizer (scan speed 19–21 Hz). To maintain the atomizer-substrate distance, the substrate was displaced vertically during spraying at a withdrawal rate of 0.7–1.04 mm/s. In order to obtain billets of different diameters, the scanning mechanism was adjusted to increase or decrease the outer scan radius. The selection of the atomizer–substrate distance
and other processing variables used was prompted by experimental and numerical results presented elsewhere [4]. These results indicated that at this flight distance, the distribution of atomized droplets contained liquid, semisolid and solid droplets [5]. 2.2. Thermal mechanical processing and heat treatment Six alloys of various compositions were processed by spray metal forming for this investigation. Table 1 displays the chemical composition and designation of the alloys. Spray formed billets of each alloy composition were then hot extruded at 400 C into round rod at an extrusion ratio of 44:1. Alloys were then solution treated between 470 and 490 C for 0.5–2 h, cold water quenched and then peak aged to a T6 heat treatment. The extruded and heat treated rods were then cut into bars 76.2 mm (3.0000 ) in length to make sub-size tensile specimen. The bars were then turned down on a lathe to a nominal diameter of 6.35 mm (0.25000 ) in the test section. The grip diameter was 7.62 mm (0.30000 ) with threaded ends to prevent slipping while testing. A minimum fillet radius of 3/1600 was used, with a reduced section at least 31.75 mm (1.2500 ) in length. The samples were machined according to specifications stated in ASTM E8-95a, for sub-size specimens. A minimum of five samples per alloy composition were used for uniaxial tensile testing. The microstructure of the extruded and heat treated alloys was observed by means of transmission electron microscopy and small-angle neutron scattering (SANS).
3. Results and discussion
Fig. 1. A schematic cross-section of the spray metal forming process showing how atomization and consolidation are combined into one step.
When a glide dislocation encounters an array of precipitates, it bows out with a curvature, at some critical angle 0 6 uc P p, between two precipitates before it can move on. The angle uc is a measure of the strength of the obstacle. According to Brown and HamÕs analysis [6], the critical shear stress at which a dislocation can move a large distance through an array of precipitates is a function of precipitate spacing (L) on the slip plane and the critical angle (uc). Weak precipitates can be overcome with very slight bending (uc = 180), whereas
M.M. Sharma et al. / Acta Materialia 53 (2005) 2919–2924
a strong precipitate cannot be overcome unless the dislocation practically doubles back on itself (uc = 0). In the latter case, the yield stress is defined by
1=2
;
ð2Þ
where G is the shear modulus, e is the coherency strain, l is the average radius and Vf is the volume fraction of precipitates. Given this information, the coherency strain for the present spray formed alloys with respect to volume fraction of metastable precipitates was estimated. First, in order to find the strengthening attributed only to the precipitation of structural particles, the change in yield strength is needed, and is defined as, Dry ¼ robs rAsQ ;
Alloy designation
Rp (nm)
Vf
(Rp * Vf)1/2
Yield strength (MPa)
9Zn–0.0Cu–0.36Mn–0.28Cr 9Zn–1.3Cu–0.34Mn–0.22Cr 12Zn–1.1Cu–0.29Mn–0.02Cr 12Zn–1.1Cu–0.18Mn–0.15Cr 12Zn–1.1Cu–0.20Mn–0.16Cr 8Zn–1.6Cu–3.97Mn–0.04Ag
2.42 2.53 2.61 2.71 2.73 2.73
0.047 0.049 0.059 0.064 0.064 0.065
0.337 0.352 0.392 0.416 0.418 0.421
711 794 807 800 842 849
ð1Þ
and corresponds to the Orowan stress, where T is the line tension, M is the Taylor factor and b is the Burgers vector. Under a simple elastic model, the line tension is given as T = Gb2/2, where G is the shear modulus of the precipitate and M is 3.06 for the polycrystalline facecentered cubic metals. Conversely, the weak interaction between precipitate and dislocation is called a cutthrough mechanism. Several models exist which describe this behavior. The coherency strain effect has been proven to be dominant for G.P. zones and metastable precipitates in the Al–Zn based alloys [2,7–12]. As a result, the contribution to strengthening from misfitting metastable precipitates was estimated in this study. The strengthening of alloys by misfitting coherent particles occurs by virtue of the interaction between stress fields of the precipitates and the dislocation. Due to the fact that the dislocation interacts the strain field of the precipitates that physically intersect its slip plane as well as the strain fields of precipitates that do not, it is necessary to find some means of averaging the effects of all the particles in the alloy. The methods of averaging the effects of all the precipitates in the microstructure on the yield strength have been discussed by Brown and Ham [6] and Gerold and Haberkorn [8], with only the constant differing from one theory to the other. According to Gerold and Haberkorn [8] and other researchers [9,10], the yield stress is expressed by the following equation: rys ¼ 3MGe3=2 ðlV f =bÞ
Table 2 Measured structure parameters for the spray formed alloys
ð3Þ
where the as-quenched yield strength is subtracted from the observed yield strength. The change in yield strength is thus the increment due to the formation of structural precipitates only. This value can then be plotted versus the structure parameters to evaluate the relationship between strength and structure of the materials. A summary of the measured structure parameters with the corresponding yield strength value is presented in Table 2. RP is the precipitate radius and the volume fraction of the precipitates is Vf. Using this information, the yield strength of these materials can be explained as a
function of the average radius and volume fraction of the structural precipitates. As-quenched data was obtained in this study, but was not considered reliable, due to the extremely fast precipitation rate during quenching, especially in the alloys with silver addition. However, if the as-quenched data is assumed to be approximately the same for all specimens, it is possible to estimate the coherency strain. Using Eq. (2), while all values other than precipitate radius (RP) and volume fraction (Vf) remain constant, the yield strength versus the square-root of the precipitate radius and volume fraction of precipitate was plotted, Fig. 2. The slope of the line is the fitted result of Eq. (2). The coherency strain is thus estimated using G = 26.2 GPa [13], b = 0.29 nm [6], M = 3.06 [7] and the slope of the line. Furthermore, the intercept from the equation of the slope of the line can be used as an approximation for the as-quenched yield strength. Table 3 presents the experimental as-quenched data for the spray formed alloys. Higher values obtained for the asquenched yield strength compared to the calculated value is attributed to the fast precipitation rate in these alloys. The time between solution treatment temperature and quench, even though very short, can allow for the beginning of precipitation of structural particles. This can make 900
y = 1244.3x + 319.94 850 Yield Strength (MPa)
rOrowan ¼ 1.6MT =bL
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800 750 700 650 600 0.250
0.300
0.350 Rp(Vf)
0.400 1/2
0.450
0.500
1/2
(nm )
Fig. 2. Change in the observed yield strength as a function of the average radius and volume fraction of g 0 precipitates. The slope can be described by the term of precipitation hardening.
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Table 3 As-quenched yield strength data for the spray formed alloys Alloy designation
As-quenched yield strength (MPa)
9Zn–0.0Cu–0.36Mn–0.28Cr 9Zn–1.3Cu–0.34Mn–0.22Cr 12Zn–1.1Cu–0.29Mn–0.02Cr 12Zn–1.1Cu–0.18Mn–0.15Cr 12Zn–1.1Cu–0.20Mn–0.16Cr 8Zn–1.6Cu–3.97Mn–0.04Ag
341 350 320 335 349 379
the experimental value for as-quenched strength slightly higher than the actual value. Table 4 presents the value for estimated coherency strain of the Al–Zn–Mg–Cu spray formed alloys and for other Al–Zn based alloy systems. From this estimation, it appears that the Al–Zn–Mg–Cu alloy system has a greater coherency strain than other systems, and processing these alloys through spray forming results in a further increase in coherency strain. The estimated value for coherency strain compares well with that of P/M processed high solute alloys. However, the data from that study was for high manganese alloys, alone. As a result, a slightly higher coherency strain was realized in those alloys. Nevertheless, the estimation validates the theory of contribution to strengthening by misfitting precipitates. Gerold and Haberkorn [8] established that after the precipitate exceeds a certain size r, the increment in flow-stress, s, with further aging increases with volume fraction of precipitate in proportion with (Vf)1/2. They explained this behavior with the help of data from Dash and Fine [11]. It was considered that Vf was essentially constant after a specific time of aging. However, the flow-stress continues to increase with aging time, despite the fact that Vf remains constant. During aging, as the particle coarsens, the average particle radius becomes larger obeying the power law r = f(t)1/3, while the volume fraction remains constant. Then, the structure parameters can be explained in terms of the Orowan stress expressed in Eq. (1), rewritten as rOrowan ¼ ð0.553MGbðV f Þ1=2 Þ=r;
ð4Þ
where r is the precipitate radius. In this equation, the Orowan stress [6,14] will decrease with increasing precipitate radius, while the yield stress from the cutthrough model will continue to increase (Eq. (2)). Using these equations, Osamura et al. [12] proposed that the theoretical maximum yield strength of these alloys would be found to be at the cross-over point for both
Fig. 3. Bright field (a) and dark field (b) electron photomicrograph of massive structural precipitation in the 12Zn–1.1Cu–0.20Mn–0.16Cr spray formed alloys.
these mechanisms. It was also proposed that the yield strength would increase with increasing coherency strain and volume fraction of precipitate, and that the strength of 1 GPa could theoretically be achieved in high solute
Table 4 Coherency strain (e) for several Al–Zn based alloy systems Al–Zn
Al–Zn–Mg
Al–Zn–Mg–Cu
Al–Zn–Mg–Cu
Al–Zn–Mg–Cu
[8] 0.007
[2] 0.015
Commercial [10] 0.022
P/M-high solute [8] 0.031
Spray formed-high solute 0.029
M.M. Sharma et al. / Acta Materialia 53 (2005) 2919–2924
concentration alloys. The solution for the theoretical yield stress at this cross-over point according to Osamura et al. [12] is as follows: rys;max ¼ 1.71MGeðV f Þ1=2 :
ð5Þ
It is clear from this analysis that three structure parameters, coherency strain, volume fraction and shear modulus, are important to the strengthening of high solute alloys. In this study, the spray formed alloys contained a variety of solute elements and an increased amount of zinc content compared to conventional 7xxx series alloys. As a result, a high volume fraction of metastable precipitates resulted in the spray formed materials and are shown in Fig. 3. The variations of amounts of zinc, magnesium, copper and silver content contributed to resulting microstructure. The estimation of the high coherency strain achieved in the spray formed alloys can be explained as follows. Chromium or manganese additions promote the formation of the rod-like E-phase or the quasi-ternary (Al20(Cu,Zn)2Mn3) intermediate phases, respectively, thus providing a fiber-like reinforcement Fig. 4. The zinc is distributed in both the E- and quasi-ternary phases, in addition to the g 0 phase. However, the magnesium is distributed in great amount in only the g 0 phase, while the E-phase has some distribution of magnesium and the quasi-ternary phase contains none. The high magnesium content tends to generate a high coherency strain due to the atomic size difference with aluminum. The microstructural characteristics of the alloys can be used to rationalize the observed mechanical properties. The estimation of coherency strain was performed in this work to correlate the increase in yield strength
Fig. 4. Bright-field photomicrographs of the fiber-like dispersoid Ephase in the 9Zn–0.0Cu–0.36Mn–0.28Cr alloy.
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Table 5 Measured structural particle concentration for various compositions Alloy designation
Particle concentration
9Zn–0.0Cu–0.36Mn–0.28Cr 9Zn–1.3Cu–0.34Mn–0.22Cr 12Zn–1.1Cu–0.29Mn–0.02Cr 12Zn–1.1Cu–0.18Mn–0.15Cr 12Zn–1.1Cu–0.20Mn–0.16Cr 8Zn–1.6Cu–3.97Mn–0.04Ag
3.0 · 104 3.5 · 104 4.0 · 104 4.0 · 104 4.0 · 104 4.5 · 104
T6
with the dependence of the size and volume fraction of metastable precipitates. This is attributed to the increasing zinc content of the material. For the most part, with similar processing conditions, increasing the zinc content resulted in an increased volume fraction of metastable precipitates, Table 5. Considering the high coherency strain estimated for the high solute spray formed alloys, and the measured high volume fraction of metastable precipitates, it can be hypothesized that the increased addition of zinc, and thus a high volume fraction of structural particles, results in an increased coherency strain. However, to verify this theory, the actual coherency strain must be measured. Measuring the coherency strain of fine precipitates cannot effectively be measured using transmission electron microscopy, since the misfit strain is often below 1% and the error is approximately 5%.
4. Summary The characteristic microstructure of the spray formed alloys is predominantly associated with a massive precipitation of g 0 plates. This microstructure is most easily achieved with alloys containing a high amount of solute content, with Zn + Mg contents over 15 wt.% where the zinc content is at least 12 wt.%. A massive and fine distribution of the g 0 phase can also be achieved with lower amounts of Zn and Mg, if accompanied by the addition of manganese and minor amounts of silver (less than 0.04 wt.%). In this study, the addition of manganese precipitated the rod-like quasi-ternary phase, which provided additional strengthening through fiber reinforcement. Therefore, it was concluded the ultrahigh strengths of the spray formed alloys is related to multiple effects of both precipitation hardening and fiber reinforcement. The reason for increased yield strength contributed from only precipitation hardening in the spray formed alloys can be explained in terms of coherency strain and a coherency strain model. The coherency strain was estimated using a model and found to be much higher in the spray formed alloys compared to PM or commercial alloys of similar composition.
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