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Cryomilled aluminum alloy and boron carbide nano-composite plate
Journal of Materials Processing Technology 209 (2009) 5046–5053
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Cryomilled aluminum alloy and boron carbide nano-composite plate R.G. Vogt, Z. Zhang, T.D. Topping, E.J. Lavernia, J.M. Schoenung ∗ Department of Chemical Engineering and Materials Science, University of California, One Shields Ave, Davis, CA 95616, USA
a r t i c l e
i n f o
Article history: Received 17 October 2008 Received in revised form 2 February 2009 Accepted 6 February 2009 Keywords: Cryomilling MMC Hot isostatic pressing Extrusion Quasi-isostatic forging
a b s t r a c t The addition of ceramic particulate reinforcement via cryomilling can significantly increase the physical and mechanical properties of Al alloys. In the present study, boron carbide (B4 C) was cryomilled with Al 5083 to form a nano-grained metal matrix powder. This powder was blended with unmilled Al 5083 to increase ductility and was then consolidated into plates by three methods: (1) hot isostatic pressing (HIPping) followed by high strain rate forging (HSRF), (2) HIPping followed by two-step quasi-isostatic forging (QIF), and (3) three-step QIF. The effects of process method on microstructure and mechanical behavior for the final consolidated nano-composite plates were investigated. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Particulate reinforced aluminum alloy metal matrix composites (MMCs) processed by mechanical milling have been reported to exhibit increases in strength, fracture toughness and stiffness over conventional aluminum alloys, which render them attractive candidates for various structural applications (Ye et al., 2005, 2006a,b). Mechanical milling produces a nano-structured (NS) matrix with homogeneously distributed reinforcement (Fogagnolo et al., 2002) and the reinforced aluminum alloy powders are subsequently processed through a powder metallurgy procedure to form bulk materials. It has been reported that significant improvement in mechanical properties has been achieved in aluminum alloys reinforced with various reinforcements such as AlN (Fogagnolo et al., 2006), B4 C (Ye et al., 2006a,b), SiC (Tang et al., 2005) and Si3 N4 (Fogagnolo et al., 2002). Previous studies of cryomilled Al-B4 C nano-composites with unmilled powder additions have shown that this material exhibits remarkable combinations of high strength and ductility (Ye et al., 2005; Witkin et al., 2003). The mechanisms reportedly responsible for the simultaneously improved strength and ductility in these multi-phase materials involves increased dislocation activity in the coarse-grained (i.e., unmilled powder) regions as a result of the constraint of plastic deformation in these coarse-grained regions (Hayes et al., 2001). Nonetheless, the reinforced nano-composites were processed via extrusion into cylindrical geometries (i.e., rods and bars). In many engineering applications, however, sheets and plates are required. Thus it is of interest to study the effects of dif-
∗ Corresponding author. E-mail address: [email protected] (J.M. Schoenung). 0924-0136/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.jmatprotec.2009.02.002
ferent processing techniques (e.g., forging) on the microstructural evolution and mechanical behavior of Al-B4 C nano-composites. B4 C as a particulate reinforcement has several advantages due to its combination of ultra-high hardness and low density; in fact this material has a density which is less than that of Al 5083 at 2.51 g/cm3 (Ye et al., 2005). Al 5083 is widely used in military applications where high strength, weldability, and corrosion resistance are necessary constraints (Newbery et al., 2007). Moreover, in related studies, cryomilled Al 5083 reportedly exhibited good thermal stability, which was generally rationalized by grain boundary segregation, grain boundary drag by solute elements, and second phase pinning, i.e., due to the formation of a small amount of nanoscale aluminum oxide, nitride and carbide dispersions during the cryomilling process (Witkin and Lavernia, 2006). In view of the above findings, Al 5083-plus-B4 C nanocomposites represent an attractive candidate for scale-up production manufacturing of plates and other industrially significant geometries. In this study, three cryomilled Al 5083-plus-B4 C nanocomposite plates were produced using three different consolidation methods: (1) hot isostatic pressing (HIPping) followed by high strain rate forging, (2) HIPping followed by two-step quasi-isostatic forging (QIF), and (3) three-step QIF. The microstructural and mechanical performance that results from each consolidation method were characterized in detail and compared. 2. Experimental procedure 2.1. Feedstock powder The feedstock powders used in this study were gas-atomized Al 5083 and Tetrabor© B4 C. The Al 5083 powder had an average particle size of <45 m and was provided by Valimet, Inc. (Stockton, CA).
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The B4 C (F1200) powder was purchased from ESK Ceramics (Saline, MI), with a particle size range of 1–7 m. Prior to cryomilling, the feedstock powders were then blended in a V-shape blender to produce a powder mixture consisting of 80 wt. pct. Al 5083 powder and 20 wt. pct. B4 C particulate. 2.2. Cryomilling Cryomilling was carried out in six sequential experiments using a 1S Szegvari attritor. For each experiment, 1 kg of blended powder (80 wt. pct. Al 5083 + 20 wt. pct. B4 C particulate) was cryomilled for 8 h. The balls and tank were immersed with liquid nitrogen during the intervals between the sequential runs while the cryomilled powder was discharged and collected. The milling media consisted of stainless steel balls submersed in liquid nitrogen with a ball-topowder ratio of 32:1 and an attritor speed of 180 rpm. Additions of 0.4 wt. pct. stearic acid were added to each batch to prevent excessive agglomeration of powder build-up on the milling balls. Further information on the details of the cryomilling technique can be found in the references by Ye et al. (2006a,b), Witkin and Lavernia (2006) and Tellkamp and Lavernia (1999). 2.3. Degassing The cryomilled powder was V-blended with an equal wt. pct. of unmilled Al 5083 powder to form a composite consisting of 10 wt. pct. B4 C, 40 wt. pct. nanocrystalline (NC) Al 5083 and 50 wt. pct. unmilled Al 5083. While contained in a nitrogen glove box, the composite powders were packed into Al 6061 cans measuring 10.16 cm in diameter by 15.24 cm tall that had a 0.95 mm stem with a valve that allowed connection to a vacuum degassing system without exposure to atmosphere. All of the canned powders were hot vacuum degassed at 450 ◦ C for 22 h after utilizing a 46-h ramp procedure. During degassing, as the temperature is increased, H2 O and H2 are removed from the powder surface oxide layer by evaporating the physisorbed H2 O and decomposing hydroxides (Young-Won et al., 1985) introduced by stearic acid, used as a process control agent (Newbery et al., 2007). After degassing, the Al 6061 cans were weld sealed to prevent atmospheric exposure prior to subsequent consolidation. 2.3.1. Hot isostatic pressing and high strain rate forging An Al 6061 can containing 1026 grams of the degassed composite powder was hot isostatic pressed (HIPped) at Kittyhawk, Inc. (Garden Grove, CA) at 400 ◦ C using a pressure of 103 MPa for 4 h. After HIPping, the Al 6061 was machined off leaving a cylindrical billet 7.26 cm in diameter and 7.24 cm tall that weighed 802.4 g. After degassing and primary consolidation, secondary processing of powder metallurgy processed aluminum composites is necessary to fully consolidate the material and break up prior particle boundary layers (Lloyd, 1994). The HIPped billet was thus high strain rate forged (HSRF) at Pittsburgh Materials Technology Inc. (Pittsburgh, PA) using a Dynapak press. The billet was forged into a die 12.7 cm square by 1.9 cm thick, using a billet pre-heat of 500 ◦ C. The pre-heating temperature was selected on the basis of secondary processing temperatures used previously to extrude cryomilled Al 5083 (Han et al., 2005). After forging, the resulting plate cracked at the corners. The material containing the cracks was removed before subsequent rolling. This plate is hereafter referred to as the HIP/HSRF plate and is shown in Fig. 1a. 2.3.2. Hot isostatic pressing and two-step quasi-isostatic forging An Al 6061 can containing 1007 g of the degassed composite powder was HIPped at Kittyhawk, Inc. (Garden Grove, CA) using the same HIPping conditions as stated above. The final consolidated cylindrical billet after machining weighed 869.4 g and was 7.45 cm
Fig. 1. Digital images of (a) HIP/HSRF plate, (b) HIP/QIF plate and (c) HIP/QIF plate. Material is Al 5083-plus-B4 C nano-composite.
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in diameter and 7.57 cm tall. The HIPped billet was quasi-isostatic forged (QIF) at Advanced Materials & Manufacturing Technologies, LLC (Roseville, CA), where it underwent two separate forgings. Due to the limitations of the allowable deformation during QIF, the as HIPped billet was first forged to 4.8 cm and then forged a second time to 3.25 cm and a diameter of 12.95 cm. Both forgings used a pre-heat temperature of 450 ◦ C, which is typically used to QIF cryomilled Al 5083 (Newbery et al., 2007). After machining, in preparation for rolling, the billet was now 2.27 cm thick and weighed 742 g, 73.6 pct. of the original powder weight. This plate is hereafter referred to as the HIP/QIF plate and is shown in Fig. 1b. 2.3.3. Three-step quasi-isostatic forging An Al 6061 containing 1030 g of degassed composite powder was QIF three times at Advanced Materials & Manufacturing Technologies, LLC (Roseville, CA). To study the influence of the primary consolidation step on the mechanical behavior of the composite, QIF was employed to consolidate the powder. Prior to the initial forging, the can was subjected to a pre-heat of 450 ◦ C for 3 h. The subsequent forgings were subjected to a pre-heat of 450 ◦ C for 2 h and then 45 min, respectively. The initial forging consolidated the composite powder to dimensions after machining of 9.65 cm in diameter and 3.99 cm tall with a weight of 785 g. The second and third forgings reduced the height by 24 and 18 pct., respectively. Prior to rolling, the final dimensions of the machined billet were 2.16 cm thick and 12.7 cm in diameter weighing 712 g, 69.1 pct. of the original powder weight. This plate is hereafter referred to as the QIF plate and is shown in Fig. 1c. 2.4. Rolling After forging, all three plates were rolled at Pittsburgh Materials Technology Inc. (Pittsburgh, PA) to a final thickness of 1.27 cm. The plates were pre-heated to 500 ◦ C for 2 h and underwent rolling at a 10 pct. reduction in thickness and 15 min re-heats during each pass to the desired thickness. 2.5. Standard Al 5083 H131 plate Standard Al 5083 H131 ‘armor grade’ plate was compared against the composite plates. Processing details include standard ingot metallurgy followed by successive hot rolling and cold working to attain the H131 condition (Newbery et al., 2007). 2.6. Microstructure characterization Chemical analysis of powders and bulk materials was completed by Luvak, Inc. (Boylston, MA). Measurement of non-metallic elements was conducted using combustion infrared detection and inert gas fusion according to ASTM E 1019-03, ASTM E 1019-3 and ASTM E 1447-05. Measurement of B4 C was conducted by gravimetric analysis by dissolving away the Al 5083 (HCl:HNO3 insoluble) from the B4 C. For metallic elements, direct current plasma emission spectroscopy according to ASTM E 197-03 was used to determine chemical composition. Forged microstructures were examined using transmission electron microscopy (TEM) and scanning electron microscopy (SEM) in both the forging and perpendicular to forging directions. TEM samples were prepared by jet polishing and mechanical thinning followed by ion-milling. Grain size measurements from TEM images in the forging and perpendicular to the forging directions were conducted to determine grain size and aspect ratio. Image analysis was conducted on the SEM images using Analysis*. X-ray diffraction (XRD) was employed to measure the average grain size of the cryomilled powder using a SINTAG X-ray diffractometer with Cu K␣ radiation. Average grain
size for the cryomilled and degassed powders before consolidation were determined from peak broadening in the X-ray diffraction pattern (Klug and Alexander, 1974). 2.7. Mechanical behavior Microhardness values were determined from the average of five indentations using a Buehler Micromet 2004 Vickers indentor. Microhardness measurements were taken in the nanocrystalline (NC) and coarse-grained regions (CG), respectively, using a 25 g load to only indent each respective region. Using a 500 g load, the microhardness indentations encompassed the CG and NC regions. Tensile specimens were Electrical Discharge Machined (EDM) to flat sub-size tensile specimens with a gauge section (40 mm × 6 mm × 3 mm) perpendicular to the forging direction according to ASTM E8. Compression specimens were also machined using EDM into 5 mm edge length cubes. Tensile and compression testing was carried out using a strain rate of 10−3 s−1 . Standard Al 5083 H131 tensile data is referenced from (Horn, 1967). 3. Results 3.1. Chemical analysis All three plates were consolidated from the same batch of cryomilled powder. The chemical composition of both feedstock powders and the cryomilled/blended powder are shown in Table 1. The bulk plates exhibit similar compositions of impurity elements, as shown in Table 1. Compared to Standard Al 5083 H131 plate (Newbery et al., 2007), the composite plates showed higher levels of iron, carbon, oxygen, hydrogen and nitrogen impurity levels, as expected after cryomilling (Witkin and Lavernia, 2006). 3.2. Microstructure The HIP/HSRF, HIP/QIF and QIF plates exhibited similar microstructures, as shown in the SEM micrographs in Fig. 2. All three plates showed lighter CG bands that do not contain B4 C particulates interwoven with darker NC bands embedded with B4 C, when looking perpendicular to the forging direction. Average grain size for the Al 5083 matrix in the cryomilled powder before blending and consolidation, as determined by XRD, showed a starting grain size of 25 nm. A small portion of cryomilled powder was degassed without blending to evaluate grain growth during degassing and showed an average grain size of 168 nm in the matrix, as determined by XRD. TEM images from the NC regions in the as-consolidated plates in the longitudinal axis direction (perpendicular to the forging direction) are shown in Fig. 3. Grain size measurement histograms from TEM images for each processing condition and orientation, i.e., in the transverse axis direction (parallel to the forging direction) and longitudinal axis direction are shown in Fig. 4. The HIP/HSRF plate had an average grain size in the NC region of 210 nm and an aspect ratio of 1.9. The HIP/QIF and QIF plates exhibited average grain sizes of 208 and 181 nm, respectively, in the longitudinal axis direction with aspect ratios of 1.6 and 1.3. 3.3. Mechanical behavior The microhardness in the NC regions of the HIP/HSRF plate has an average value of 232 HV, which is slightly higher than that of the HIP/QIF and QIF plates (i.e., 214 and 216 HV, respectively), as shown in Table 2. The microhardness in the CG regions of the HIP/QIF and QIF plates is slightly higher at 99 and 100 HV, respectively, than for the HIP/HSRF plate with a value of 85 HV. Hardness for the Standard Al 5083 H131 plate was 20 pct. and 10 pct. lower than the HIP/HSRF and QIF composite plates, respectively. Indentations made with the
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Table 1 Chemical analysis of powders and Al 5083-plus-B4 C nano-composite plates. Values in wt. pct., except H (ppm). Condition
Processing
C
O
Powders
Al 5083 B4 C Cryomilled and blended
0.009 22 2.04
0.18 0.50 1.03
Plates
HIP/HSRF HIP/QIF QIF H131a
a
2.2 2.1 2.1 0.001
H
0.46 0.54 0.46 0.003
1700 350 380 62 44 49 1.3
N
Fe
B4 C
0.002 0.08 0.51
0.28 0.49 0.26
– – 9.25
0.33 0.24 0.32 <0.005
0.24 0.29 0.21 0.4
– – – –
Chemical analysis data for Standard Al 5803 H131 referenced from (Newbery et al., 2007).
500 g load show that the overall hardness value (accounting for both NC/CG regions) is 150 HV for the HIP/HSRF plate, which is higher than that of the HIP/QIF forged plate, as shown in Table 2. The QIF plate exhibits a slight increase in average hardness over the HIP/HSRF plate.
In tension, the yield strength at 0.2 pct. offset for the HIP/QIF and QIF plates was 320 and 318 MPa with elongations of 2.15 pct. and 2.0 pct. respectively, as shown in Table 2. The HIP/HSRF plate showed higher yield strength at 397 MPa coupled with a lower elongation value of 0.9 pct. The QIF plates yielded at values 20 pct. higher than
Fig. 2. SEM micrographs of (a) HIP/HSRF plate perpendicular to forging direction, (b) HIP/HSRF plate parallel to forging direction, (c) HIP/QIF plate perpendicular to forging direction, (d) HIP/QIF plate parallel to forging direction, (e) QIF plate perpendicular to forging direction, and (f) QIF plate parallel to forging direction. Material is Al 5083-plus-B4 C nano-composite.
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R.G. Vogt et al. / Journal of Materials Processing Technology 209 (2009) 5046–5053 Table 3 Compression properties of the Al 5083-plus-B4 C nano-composite plates. Data on standard Al 5083 H131 is provided for comparison. Processing
Orientation
0.2% yield strength (MPa)
Ultimate strength (MPa)
Compressive strain (%)
HIP/HSRF
Longitudinal Transverse
473 464
636 618
7.2 5.2
HIP/QIF
Longitudinal Transverse
386 358
609 566
22.3 12.9
QIF
Longitudinal Transverse
398 348
606 572
20.4 11.6
H131a
Longitudinal Transverse
288 266
426 421
>60 >60
a Compression data on Standard Al 5083 H131 was not found in the literature. The values presented here were derived from tests conducted in our laboratory. Results represent one compression test per direction with 8 mm edge length cubes. Samples were compressed to 50 kN, after which the test was stopped to avoid damage to the Instron test apparatus, as samples did not fracture.
a ductile fracture morphology corresponding to the unreinforced CG regions. The darker interparticle reinforced NC regions exhibited a brittle fracture morphology. Compression testing on the plates was conducted in both the forging direction and perpendicular to the forging direction. The plates exhibited more ductility and higher strength in the direction perpendicular to the forging, as shown in Table 3. The yield strength at 0.2 pct. offset for the HIP/QIF and QIF plates were 358 and 348 MPa, respectively, in the forging direction and 386 and 398 MPa, respectively, perpendicular to the forging direction. The HIP/HSRF plate yielded at a value nearly 100 MPa higher than the QIF plates in both the perpendicular and parallel to the forging directions. As a consequence of increased strength, the HIP/HSRF plate had significantly less ductility than the QIF plates, as shown in Table 3. In the longitudinal direction, the HIP/HSRF plate and QIF plates showed a 40 pct. and 25 pct. increase in yield strength respectively over the Standard H131 plate. 4. Discussion
Fig. 3. TEM images of (a) HIP/HSRF plate, (b) HIP/QIF plate, and (c) QIF plate. Material is Al 5083-plus-B4 C nano-composite.
the Standard H131 plate, and the HIP/HSRF plate yielded at a value 35 pct. higher, both in the longitudinal direction. Fig. 5 shows the fracture surfaces after tension testing for the HIP/HSRF, HIP/QIF and QIF samples. The SEM micrographs of the fracture surfaces indicate
Chemical analysis on the as received powders falls within manufacturer specifications for both Al 5083 and B4 C. For the cryomilled-and-blended powder there is an increase in the carbon, oxygen, nitrogen, hydrogen and iron content, as shown in Table 1, compared to Standard Al 5083 H131. Despite the handling of powders under the cover of liquid nitrogen after cryomilling and keeping the powders in a protected argon atmosphere, the surface of aluminum is prone to forming a thin oxide layer from oxygen in the environment that reacts with the powder during processing; extremely low O2 partial pressure, <10−143 torr, is necessary to prevent oxidation (Kowalski et al., 1992). Increased hydrogen contamination results from hydroxides and stearic acid, which coupled with oxygen contamination can have a detrimental effect on ductility (Lu et al., 2002). A characteristic of cryomilled materials is the
Table 2 Tensile and hardness properties of the Al 5083-plus-B4 C nano-composite plates in the longitudinal orientation. Data on Standard Al 5083 H131 are provided for comparison. Processing
0.2% yield strength (MPa)
Ultimate tensile strength (MPa)
Elongation (%)
Elastic modulus (GPa)
HIP/HSRF HIP/QIF QIF H131a
397 320 318 255
502 471 460 338
0.9 2.0 1.7 14.0
80.5 82 81 71
NC, Nanocrystalline region; CG, coarse-grained region. a Standard Al 5083 H131 referenced from (Horn, 1967).
Hardness (Vickers) NC/CG
NC
150 ± 11 136 ± 9 158 ± 9 –
232 ± 22 214 ± 18 216 ± 16 –
CG 85 ± 7 99 ± 3 100 ± 4 90
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Fig. 4. Grain size distribution histograms for: (a) HIP/HSRF longitudinal axis, (b) HIP/HSRF transverse axis, (c) HIP/QIF longitudinal axis, (d) HIP/QIF transverse axis, (e) QIF longitudinal axis, and (f) QIF transverse axis. Material is Al 5083-plus-B4 C nano-composite.
increased thermal stability due to small dispersoids that result from the presence of the oxygen and nitrogen, giving rise to Zener pinning (Witkin and Lavernia, 2006). The stearic acid used as a process control agent is the primary source of the carbon and hydrogen in the cryomilled-and-blended powder; removal of these is the primary purpose of the degassing step. The hydrogen content decreased ∼330 ppm from the cryomilled-and-blended powder to the final consolidated composite plates; this reduction can be attributed to the degassing step. As a trade-off, however, the degassing also results in grain coarsening of the Al 5083 matrix from ∼25 nm after cryomilling to ∼168 nm after degassing. The mechanical behavior of the composite plates in tension and compression exhibited an asymmetry in yield strength while the Standard H131 plate exhibited some symmetry. The HIP/HSRF plate yielded at a value ∼100 MPa higher in compression than in tension. The QIF plates yielded roughly 60 and 70 MPa higher in compression than in tension for the HIP/QIF and QIF plates, respectively. This asymmetry behavior has been observed before by Ye et al. (2006a,b) in extruded MMC’s and is a generic feature in composites (Clyne
and Withers, 1993). Upon cooling from the processing temperature there is a residual tensile stress field that remains in the composite (Fernández et al., 2004) resulting in a lower tensile yield strength. The compression testing in the forging direction and perpendicular to the forging direction for all three plates show anisotropic behavior that can be explained by the anisotropic microstructure that results during deformation (Skolnik et al., 2008). The fracture surfaces, shown in Fig. 5, exhibit well-defined dimples in the CG bands indicating a ductile fracture mechanism after all three processing methods. In the particulate reinforced NC matrix regions, the fracture morphology is not as distinct as in the unreinforced CG regions with much shallower peaks and valleys. B4 C particles are not apparent at the NC matrix fracture surfaces suggesting that debonding between the matrix and reinforcement is not a predominant fracture mechanism, which is consistent with the behavior observed previously in aluminum reinforced with SiC particulate (Beffort et al., 2007). The HIP/HSRF plate showed higher strength values, with limited ductility, than the plates processed by QIF. The grain size distri-
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bution data shown in Fig. 4 indicate that the distributions in the NC regions are similar in the three plates, and therefore, grain size strengthening does not appear to be the reason for the observed differences in mechanical behavior. Considering the consolidation techniques in more detail, it is important to note that HSRF operates at an extremely high strain rate, 102 s−1 , compared to QIF at 10−2 s−1 . This high strain rate can prevent dynamic recovery or recrystallization during the processing; this is supported by the evidence that the HSRF plates showed a higher aspect ratio (1.9) compared to that of the QIF plates (1.6 and 1.3). The lower aspect ratio in the QIF plates reflects a large portion of recrystallized grains in the QIF plates. As shown in Fig. 4, the grain size distribution
Fig. 6. Work hardening rate for the three plates consolidated using three different methods based on compression results in the transverse and longitudinal orientation. Material is Al 5083-plus-B4 C nano-composite.
widened in the transverse direction and narrowed in the longitudinal direction in the QIF forged plates compared to that in the HSRF plates. To further consider this effect, the work hardening rate can be calculated by =
Fig. 5. Tensile fracture surfaces of (a) HIP/HSRF plate, (b) HIP/QIF plate, and (c) QIF plate. Material is Al 5083-plus-B4 C nano-composite.
d dε
(1)
where is the change in stress divided by the change in strain (Rodriguez et al., 2003), the results of which are shown in Fig. 6 calculated using true stress–strain compression data. The work hardening rate is higher for the QIF plates than for the HSRF plate, which is in agreement with the lower strengths and increased ductility of the QIF plates, suggesting that the QIF plates have lower initial dislocation density. In nanocrystalline materials, recovery or recrystallization may occur via rotation or coalescence of adjacent smaller grains (Zhou et al., 2003). The extent of grain rotation and coalescence is expected to increase during deformation in response to the orientation accommodation between adjacent grains under an applied load (Jin et al., 2004) or driven by stress-assisted grain boundary migration (Jin et al., 2004; Gianola et al., 2006). In a study by Witkin et al. (2005), evaluating microstructural coarsening during thermomechanical processing, they found that recovery or recrystallization occurred quickly after the onset of deformation but was not sensitive to the strain. Therefore, in addition to the strain rate effect, the larger extent of recrystallization in the QIF plates can also be caused by the fact that the QIF plates underwent more deformation steps than the HSRF plates. This view suggests that a single pass high strain rate consolidation method is preferred to achieve high strength while subsequent thermomechanical processing can be used to increase the ductility in the Al 5083-plus-B4 C nano-composite.
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5. Conclusions Cryomilled Al 5083-plus-B4 C powder blended with unmilled CG Al 5083 powder was successfully consolidated into plates via three methods: (1) hot isostatic press (HIP) and high strain rate forge (HSRF), (2) HIP and two-step quasi-isostatic forge (QIF), and (3) three-step QIF, producing three Plates 1.3 cm thick and approximately 15 cm in diameter. The HIP/HSRF plate exhibited higher strength with less ductility than the QIF plates, which had similar mechanical properties. The increased strength and reduced ductility of the HIP/HSRF plate is attributed to the inhibition of dynamic recrystallization during the high strain rate forging procedure. Acknowledgements Financial support for this work was provided by the Office of Naval Research, through contract N00014-03-C-0163 and U.S. Navy STTR Phase II Contract No. N00167-08-C-0001. References Beffort, O., Long, S., Cayron, C., Kuebler, J., Buffat, P.-A., 2007. Alloying effects on microstructure and mechanical properties of high volume fraction SiC-particle reinforced Al-MMCs made by squeeze casting infiltration. Composites Science and Technology 67, 737–745. Clyne, T.W., Withers, P.J. (Eds.), 1993. An Introduction to Metal Matrix Composites. Press Syndicate of the University of Cambridge, Cambridge, pp. 71–112. Fernández, R., Bruno, G., González-Doncel, G., 2004. Correlation between residual stresses and the strength differential effect in PM 6061Al-15 vol% SiCw composites: experiments, models and predictions. Acta Materialia 52, 5471–5483. Fogagnolo, J.B., Robert, M.H., Torralba, J.M., 2006. Mechanically alloyed AlN particlereinforced Al-6061 matrix composites: powder processing, consolidation and mechanical strength and hardness of the as-extruded materials. Materials Science and Engineering: A 426, 85–94. Fogagnolo, J.B., Ruiz-Navas, E.M., Robert, M.H., Torralba, J.M., 2002. 6061 Al reinforced with silicon nitride particles processed by mechanical milling. Scripta Materialia 47, 243–248. Gianola, D.S., Van Petegem, S., Legros, M., Brandstetter, S., Van Swygenhoven, H., Hemker, K.J., 2006. Stress-assisted discontinuous grain growth and its effect on the deformation behavior of nanocrystalline aluminum thin films. Acta Materialia 54, 2253–2263. Han, B.Q., Lee, Z., Witkin, D., Nutt, S., Lavernia, E.J., 2005. Deformation behavior of bimodal nanostructured 5083 Al alloys. Metallurgical and Materials Transactions A—Physical Metallurgy and Materials Science 36A, 957–965. Hayes, R.W., Rodriguez, R., Lavernia, E.J., 2001. The mechanical behavior of a cryomilled Al–10Ti–2Cu alloy. Acta Materialia 49, 4055–4068.
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Horn, K.R.V. (Ed.), 1967. Aluminum. Design and Application. Chapman & Hall, Metals Park, Ohio, pp. 175–179. Jin, M., Minor, A.M., Stach, E.A., Morris, J.J.W., 2004. Direct observation of deformation-induced grain growth during the nanoindentation of ultrafinegrained Al at room temperature. Acta Materialia 52, 5381–5387. Klug, H.P., Alexander, L.E., 1974. X-Ray Diffraction Procedures. John Wiley & Sons, New York, pp. 618–705. Kowalski, L., Korevaar, B.M., Duszczyk, J., 1992. Some new aspects of the theory of oxidation and degassing of aluminum-based alloy powders. Journal of Materials Science 27, 2770–2780. Lloyd, D.J., 1994. Particle-reinforced aluminum and magnesium matrix composites. International Materials Reviews 39, 1–23. Lu, G., Orlikowski, D., Park, I., Politano, O., Kaxiras, E., 2002. Energetics of hydrogen impurities in aluminum and their effect on mechanical properties. Physical Review B 65, 064102. Newbery, A.P., Ahn, B., Topping, T.D., Pao, P.S., Nutt, S.R., Lavernia, E.J., 2007. Large UFG Al alloy plates from cryomilling. Journal of Materials Processing Technology 203, 37–45. Rodriguez, R., Hayes, R.W., Berbon, P.B., Lavernia, E.J., 2003. Tensile and creep behavior of cryomilled Inco 625. Acta Materialia 51, 911–929. Skolnik, D.A., Liu, H.T., Wu, H.C., Sun, L.Z., 2008. Anisotropic elastoplastic and damage behavior of SiCp/Al composite sheets. International Journal of Damage Mechanics 17, 247–272. Tang, F., Hagiwara, M., Schoenung, J.M., 2005. Microstructure and tensile properties of bulk nanostructured Al-5083/SiCp composites prepared by cryomilling. Materials Science and Engineering: A 407, 306–314. Tellkamp, V.L., Lavernia, E.J., 1999. Processing and mechanical properties of nanocrystalline 5083 Al alloy. Nanostructured Materials 12, 249–252. Witkin, D., Han, B.Q., Lavernia, E.J., 2005. Microstructural evolution of an ultrafinegrained cryomilled Al 5083 alloy during thermomechanical processing. Journal of Materials Research 20, 2117–2126. Witkin, D., Lee, Z., Rodriguez, R., Nutt, S., Lavernia, E., 2003. Al–Mg alloy engineered with bimodal grain size for high strength and increased ductility. Scripta Materialia 49, 297–302. Witkin, D.B., Lavernia, E.J., 2006. Synthesis and mechanical behavior of nanostructured materials via cryomilling. Progress in Materials Science 51, 1–60. Ye, J., Han, B.Q., Lee, Z., Ahn, B., Nutt, S.R., Schoenung, J.M., 2005. A tri-modal aluminum based composite with super-high strength. Scripta Materialia 53, 481–486. Ye, J., Han, B.Q., Schoenung, J.M., 2006a. Mechanical behaviour of an Al-matrix composite reinforced with nanocrystalline Al-coated B4C particulates. Philosophical Magazine Letters 86, 721–732. Ye, J.C., He, J.H., Schoenung, J.M., 2006b. Cryomilling for the fabrication of a particulate B4C reinforced Al nanocomposite: part I. Effects of process conditions on structure. Metallurgical and Materials Transactions A—Physical Metallurgy and Materials Science 37A, 3099–3109. Young-Won, K., Griffith, W.M., Froes, F.H., 1985. Surface oxides in P/M aluminum alloys. Journal of Metals, 27–33. Zhou, F., Liao, X.Z., Zhu, Y.T., Dallek, S., Lavernia, E.J., 2003. Microstructural evolution during recovery and recrystallization of a nanocrystalline Al–Mg alloy prepared by cryogenic ball milling. Acta Materialia 51, 2777–2791.
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Cryomilled aluminum alloy and boron carbide nano-composite plate R.G. Vogt, Z. Zhang, T.D. Topping, E.J. Lavernia, J.M. Schoenung ∗ Department of Chemical Engineering and Materials Science, University of California, One Shields Ave, Davis, CA 95616, USA
a r t i c l e
i n f o
Article history: Received 17 October 2008 Received in revised form 2 February 2009 Accepted 6 February 2009 Keywords: Cryomilling MMC Hot isostatic pressing Extrusion Quasi-isostatic forging
a b s t r a c t The addition of ceramic particulate reinforcement via cryomilling can significantly increase the physical and mechanical properties of Al alloys. In the present study, boron carbide (B4 C) was cryomilled with Al 5083 to form a nano-grained metal matrix powder. This powder was blended with unmilled Al 5083 to increase ductility and was then consolidated into plates by three methods: (1) hot isostatic pressing (HIPping) followed by high strain rate forging (HSRF), (2) HIPping followed by two-step quasi-isostatic forging (QIF), and (3) three-step QIF. The effects of process method on microstructure and mechanical behavior for the final consolidated nano-composite plates were investigated. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Particulate reinforced aluminum alloy metal matrix composites (MMCs) processed by mechanical milling have been reported to exhibit increases in strength, fracture toughness and stiffness over conventional aluminum alloys, which render them attractive candidates for various structural applications (Ye et al., 2005, 2006a,b). Mechanical milling produces a nano-structured (NS) matrix with homogeneously distributed reinforcement (Fogagnolo et al., 2002) and the reinforced aluminum alloy powders are subsequently processed through a powder metallurgy procedure to form bulk materials. It has been reported that significant improvement in mechanical properties has been achieved in aluminum alloys reinforced with various reinforcements such as AlN (Fogagnolo et al., 2006), B4 C (Ye et al., 2006a,b), SiC (Tang et al., 2005) and Si3 N4 (Fogagnolo et al., 2002). Previous studies of cryomilled Al-B4 C nano-composites with unmilled powder additions have shown that this material exhibits remarkable combinations of high strength and ductility (Ye et al., 2005; Witkin et al., 2003). The mechanisms reportedly responsible for the simultaneously improved strength and ductility in these multi-phase materials involves increased dislocation activity in the coarse-grained (i.e., unmilled powder) regions as a result of the constraint of plastic deformation in these coarse-grained regions (Hayes et al., 2001). Nonetheless, the reinforced nano-composites were processed via extrusion into cylindrical geometries (i.e., rods and bars). In many engineering applications, however, sheets and plates are required. Thus it is of interest to study the effects of dif-
∗ Corresponding author. E-mail address: [email protected] (J.M. Schoenung). 0924-0136/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.jmatprotec.2009.02.002
ferent processing techniques (e.g., forging) on the microstructural evolution and mechanical behavior of Al-B4 C nano-composites. B4 C as a particulate reinforcement has several advantages due to its combination of ultra-high hardness and low density; in fact this material has a density which is less than that of Al 5083 at 2.51 g/cm3 (Ye et al., 2005). Al 5083 is widely used in military applications where high strength, weldability, and corrosion resistance are necessary constraints (Newbery et al., 2007). Moreover, in related studies, cryomilled Al 5083 reportedly exhibited good thermal stability, which was generally rationalized by grain boundary segregation, grain boundary drag by solute elements, and second phase pinning, i.e., due to the formation of a small amount of nanoscale aluminum oxide, nitride and carbide dispersions during the cryomilling process (Witkin and Lavernia, 2006). In view of the above findings, Al 5083-plus-B4 C nanocomposites represent an attractive candidate for scale-up production manufacturing of plates and other industrially significant geometries. In this study, three cryomilled Al 5083-plus-B4 C nanocomposite plates were produced using three different consolidation methods: (1) hot isostatic pressing (HIPping) followed by high strain rate forging, (2) HIPping followed by two-step quasi-isostatic forging (QIF), and (3) three-step QIF. The microstructural and mechanical performance that results from each consolidation method were characterized in detail and compared. 2. Experimental procedure 2.1. Feedstock powder The feedstock powders used in this study were gas-atomized Al 5083 and Tetrabor© B4 C. The Al 5083 powder had an average particle size of <45 m and was provided by Valimet, Inc. (Stockton, CA).
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The B4 C (F1200) powder was purchased from ESK Ceramics (Saline, MI), with a particle size range of 1–7 m. Prior to cryomilling, the feedstock powders were then blended in a V-shape blender to produce a powder mixture consisting of 80 wt. pct. Al 5083 powder and 20 wt. pct. B4 C particulate. 2.2. Cryomilling Cryomilling was carried out in six sequential experiments using a 1S Szegvari attritor. For each experiment, 1 kg of blended powder (80 wt. pct. Al 5083 + 20 wt. pct. B4 C particulate) was cryomilled for 8 h. The balls and tank were immersed with liquid nitrogen during the intervals between the sequential runs while the cryomilled powder was discharged and collected. The milling media consisted of stainless steel balls submersed in liquid nitrogen with a ball-topowder ratio of 32:1 and an attritor speed of 180 rpm. Additions of 0.4 wt. pct. stearic acid were added to each batch to prevent excessive agglomeration of powder build-up on the milling balls. Further information on the details of the cryomilling technique can be found in the references by Ye et al. (2006a,b), Witkin and Lavernia (2006) and Tellkamp and Lavernia (1999). 2.3. Degassing The cryomilled powder was V-blended with an equal wt. pct. of unmilled Al 5083 powder to form a composite consisting of 10 wt. pct. B4 C, 40 wt. pct. nanocrystalline (NC) Al 5083 and 50 wt. pct. unmilled Al 5083. While contained in a nitrogen glove box, the composite powders were packed into Al 6061 cans measuring 10.16 cm in diameter by 15.24 cm tall that had a 0.95 mm stem with a valve that allowed connection to a vacuum degassing system without exposure to atmosphere. All of the canned powders were hot vacuum degassed at 450 ◦ C for 22 h after utilizing a 46-h ramp procedure. During degassing, as the temperature is increased, H2 O and H2 are removed from the powder surface oxide layer by evaporating the physisorbed H2 O and decomposing hydroxides (Young-Won et al., 1985) introduced by stearic acid, used as a process control agent (Newbery et al., 2007). After degassing, the Al 6061 cans were weld sealed to prevent atmospheric exposure prior to subsequent consolidation. 2.3.1. Hot isostatic pressing and high strain rate forging An Al 6061 can containing 1026 grams of the degassed composite powder was hot isostatic pressed (HIPped) at Kittyhawk, Inc. (Garden Grove, CA) at 400 ◦ C using a pressure of 103 MPa for 4 h. After HIPping, the Al 6061 was machined off leaving a cylindrical billet 7.26 cm in diameter and 7.24 cm tall that weighed 802.4 g. After degassing and primary consolidation, secondary processing of powder metallurgy processed aluminum composites is necessary to fully consolidate the material and break up prior particle boundary layers (Lloyd, 1994). The HIPped billet was thus high strain rate forged (HSRF) at Pittsburgh Materials Technology Inc. (Pittsburgh, PA) using a Dynapak press. The billet was forged into a die 12.7 cm square by 1.9 cm thick, using a billet pre-heat of 500 ◦ C. The pre-heating temperature was selected on the basis of secondary processing temperatures used previously to extrude cryomilled Al 5083 (Han et al., 2005). After forging, the resulting plate cracked at the corners. The material containing the cracks was removed before subsequent rolling. This plate is hereafter referred to as the HIP/HSRF plate and is shown in Fig. 1a. 2.3.2. Hot isostatic pressing and two-step quasi-isostatic forging An Al 6061 can containing 1007 g of the degassed composite powder was HIPped at Kittyhawk, Inc. (Garden Grove, CA) using the same HIPping conditions as stated above. The final consolidated cylindrical billet after machining weighed 869.4 g and was 7.45 cm
Fig. 1. Digital images of (a) HIP/HSRF plate, (b) HIP/QIF plate and (c) HIP/QIF plate. Material is Al 5083-plus-B4 C nano-composite.
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in diameter and 7.57 cm tall. The HIPped billet was quasi-isostatic forged (QIF) at Advanced Materials & Manufacturing Technologies, LLC (Roseville, CA), where it underwent two separate forgings. Due to the limitations of the allowable deformation during QIF, the as HIPped billet was first forged to 4.8 cm and then forged a second time to 3.25 cm and a diameter of 12.95 cm. Both forgings used a pre-heat temperature of 450 ◦ C, which is typically used to QIF cryomilled Al 5083 (Newbery et al., 2007). After machining, in preparation for rolling, the billet was now 2.27 cm thick and weighed 742 g, 73.6 pct. of the original powder weight. This plate is hereafter referred to as the HIP/QIF plate and is shown in Fig. 1b. 2.3.3. Three-step quasi-isostatic forging An Al 6061 containing 1030 g of degassed composite powder was QIF three times at Advanced Materials & Manufacturing Technologies, LLC (Roseville, CA). To study the influence of the primary consolidation step on the mechanical behavior of the composite, QIF was employed to consolidate the powder. Prior to the initial forging, the can was subjected to a pre-heat of 450 ◦ C for 3 h. The subsequent forgings were subjected to a pre-heat of 450 ◦ C for 2 h and then 45 min, respectively. The initial forging consolidated the composite powder to dimensions after machining of 9.65 cm in diameter and 3.99 cm tall with a weight of 785 g. The second and third forgings reduced the height by 24 and 18 pct., respectively. Prior to rolling, the final dimensions of the machined billet were 2.16 cm thick and 12.7 cm in diameter weighing 712 g, 69.1 pct. of the original powder weight. This plate is hereafter referred to as the QIF plate and is shown in Fig. 1c. 2.4. Rolling After forging, all three plates were rolled at Pittsburgh Materials Technology Inc. (Pittsburgh, PA) to a final thickness of 1.27 cm. The plates were pre-heated to 500 ◦ C for 2 h and underwent rolling at a 10 pct. reduction in thickness and 15 min re-heats during each pass to the desired thickness. 2.5. Standard Al 5083 H131 plate Standard Al 5083 H131 ‘armor grade’ plate was compared against the composite plates. Processing details include standard ingot metallurgy followed by successive hot rolling and cold working to attain the H131 condition (Newbery et al., 2007). 2.6. Microstructure characterization Chemical analysis of powders and bulk materials was completed by Luvak, Inc. (Boylston, MA). Measurement of non-metallic elements was conducted using combustion infrared detection and inert gas fusion according to ASTM E 1019-03, ASTM E 1019-3 and ASTM E 1447-05. Measurement of B4 C was conducted by gravimetric analysis by dissolving away the Al 5083 (HCl:HNO3 insoluble) from the B4 C. For metallic elements, direct current plasma emission spectroscopy according to ASTM E 197-03 was used to determine chemical composition. Forged microstructures were examined using transmission electron microscopy (TEM) and scanning electron microscopy (SEM) in both the forging and perpendicular to forging directions. TEM samples were prepared by jet polishing and mechanical thinning followed by ion-milling. Grain size measurements from TEM images in the forging and perpendicular to the forging directions were conducted to determine grain size and aspect ratio. Image analysis was conducted on the SEM images using Analysis*. X-ray diffraction (XRD) was employed to measure the average grain size of the cryomilled powder using a SINTAG X-ray diffractometer with Cu K␣ radiation. Average grain
size for the cryomilled and degassed powders before consolidation were determined from peak broadening in the X-ray diffraction pattern (Klug and Alexander, 1974). 2.7. Mechanical behavior Microhardness values were determined from the average of five indentations using a Buehler Micromet 2004 Vickers indentor. Microhardness measurements were taken in the nanocrystalline (NC) and coarse-grained regions (CG), respectively, using a 25 g load to only indent each respective region. Using a 500 g load, the microhardness indentations encompassed the CG and NC regions. Tensile specimens were Electrical Discharge Machined (EDM) to flat sub-size tensile specimens with a gauge section (40 mm × 6 mm × 3 mm) perpendicular to the forging direction according to ASTM E8. Compression specimens were also machined using EDM into 5 mm edge length cubes. Tensile and compression testing was carried out using a strain rate of 10−3 s−1 . Standard Al 5083 H131 tensile data is referenced from (Horn, 1967). 3. Results 3.1. Chemical analysis All three plates were consolidated from the same batch of cryomilled powder. The chemical composition of both feedstock powders and the cryomilled/blended powder are shown in Table 1. The bulk plates exhibit similar compositions of impurity elements, as shown in Table 1. Compared to Standard Al 5083 H131 plate (Newbery et al., 2007), the composite plates showed higher levels of iron, carbon, oxygen, hydrogen and nitrogen impurity levels, as expected after cryomilling (Witkin and Lavernia, 2006). 3.2. Microstructure The HIP/HSRF, HIP/QIF and QIF plates exhibited similar microstructures, as shown in the SEM micrographs in Fig. 2. All three plates showed lighter CG bands that do not contain B4 C particulates interwoven with darker NC bands embedded with B4 C, when looking perpendicular to the forging direction. Average grain size for the Al 5083 matrix in the cryomilled powder before blending and consolidation, as determined by XRD, showed a starting grain size of 25 nm. A small portion of cryomilled powder was degassed without blending to evaluate grain growth during degassing and showed an average grain size of 168 nm in the matrix, as determined by XRD. TEM images from the NC regions in the as-consolidated plates in the longitudinal axis direction (perpendicular to the forging direction) are shown in Fig. 3. Grain size measurement histograms from TEM images for each processing condition and orientation, i.e., in the transverse axis direction (parallel to the forging direction) and longitudinal axis direction are shown in Fig. 4. The HIP/HSRF plate had an average grain size in the NC region of 210 nm and an aspect ratio of 1.9. The HIP/QIF and QIF plates exhibited average grain sizes of 208 and 181 nm, respectively, in the longitudinal axis direction with aspect ratios of 1.6 and 1.3. 3.3. Mechanical behavior The microhardness in the NC regions of the HIP/HSRF plate has an average value of 232 HV, which is slightly higher than that of the HIP/QIF and QIF plates (i.e., 214 and 216 HV, respectively), as shown in Table 2. The microhardness in the CG regions of the HIP/QIF and QIF plates is slightly higher at 99 and 100 HV, respectively, than for the HIP/HSRF plate with a value of 85 HV. Hardness for the Standard Al 5083 H131 plate was 20 pct. and 10 pct. lower than the HIP/HSRF and QIF composite plates, respectively. Indentations made with the
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Table 1 Chemical analysis of powders and Al 5083-plus-B4 C nano-composite plates. Values in wt. pct., except H (ppm). Condition
Processing
C
O
Powders
Al 5083 B4 C Cryomilled and blended
0.009 22 2.04
0.18 0.50 1.03
Plates
HIP/HSRF HIP/QIF QIF H131a
a
2.2 2.1 2.1 0.001
H
0.46 0.54 0.46 0.003
1700 350 380 62 44 49 1.3
N
Fe
B4 C
0.002 0.08 0.51
0.28 0.49 0.26
– – 9.25
0.33 0.24 0.32 <0.005
0.24 0.29 0.21 0.4
– – – –
Chemical analysis data for Standard Al 5803 H131 referenced from (Newbery et al., 2007).
500 g load show that the overall hardness value (accounting for both NC/CG regions) is 150 HV for the HIP/HSRF plate, which is higher than that of the HIP/QIF forged plate, as shown in Table 2. The QIF plate exhibits a slight increase in average hardness over the HIP/HSRF plate.
In tension, the yield strength at 0.2 pct. offset for the HIP/QIF and QIF plates was 320 and 318 MPa with elongations of 2.15 pct. and 2.0 pct. respectively, as shown in Table 2. The HIP/HSRF plate showed higher yield strength at 397 MPa coupled with a lower elongation value of 0.9 pct. The QIF plates yielded at values 20 pct. higher than
Fig. 2. SEM micrographs of (a) HIP/HSRF plate perpendicular to forging direction, (b) HIP/HSRF plate parallel to forging direction, (c) HIP/QIF plate perpendicular to forging direction, (d) HIP/QIF plate parallel to forging direction, (e) QIF plate perpendicular to forging direction, and (f) QIF plate parallel to forging direction. Material is Al 5083-plus-B4 C nano-composite.
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R.G. Vogt et al. / Journal of Materials Processing Technology 209 (2009) 5046–5053 Table 3 Compression properties of the Al 5083-plus-B4 C nano-composite plates. Data on standard Al 5083 H131 is provided for comparison. Processing
Orientation
0.2% yield strength (MPa)
Ultimate strength (MPa)
Compressive strain (%)
HIP/HSRF
Longitudinal Transverse
473 464
636 618
7.2 5.2
HIP/QIF
Longitudinal Transverse
386 358
609 566
22.3 12.9
QIF
Longitudinal Transverse
398 348
606 572
20.4 11.6
H131a
Longitudinal Transverse
288 266
426 421
>60 >60
a Compression data on Standard Al 5083 H131 was not found in the literature. The values presented here were derived from tests conducted in our laboratory. Results represent one compression test per direction with 8 mm edge length cubes. Samples were compressed to 50 kN, after which the test was stopped to avoid damage to the Instron test apparatus, as samples did not fracture.
a ductile fracture morphology corresponding to the unreinforced CG regions. The darker interparticle reinforced NC regions exhibited a brittle fracture morphology. Compression testing on the plates was conducted in both the forging direction and perpendicular to the forging direction. The plates exhibited more ductility and higher strength in the direction perpendicular to the forging, as shown in Table 3. The yield strength at 0.2 pct. offset for the HIP/QIF and QIF plates were 358 and 348 MPa, respectively, in the forging direction and 386 and 398 MPa, respectively, perpendicular to the forging direction. The HIP/HSRF plate yielded at a value nearly 100 MPa higher than the QIF plates in both the perpendicular and parallel to the forging directions. As a consequence of increased strength, the HIP/HSRF plate had significantly less ductility than the QIF plates, as shown in Table 3. In the longitudinal direction, the HIP/HSRF plate and QIF plates showed a 40 pct. and 25 pct. increase in yield strength respectively over the Standard H131 plate. 4. Discussion
Fig. 3. TEM images of (a) HIP/HSRF plate, (b) HIP/QIF plate, and (c) QIF plate. Material is Al 5083-plus-B4 C nano-composite.
the Standard H131 plate, and the HIP/HSRF plate yielded at a value 35 pct. higher, both in the longitudinal direction. Fig. 5 shows the fracture surfaces after tension testing for the HIP/HSRF, HIP/QIF and QIF samples. The SEM micrographs of the fracture surfaces indicate
Chemical analysis on the as received powders falls within manufacturer specifications for both Al 5083 and B4 C. For the cryomilled-and-blended powder there is an increase in the carbon, oxygen, nitrogen, hydrogen and iron content, as shown in Table 1, compared to Standard Al 5083 H131. Despite the handling of powders under the cover of liquid nitrogen after cryomilling and keeping the powders in a protected argon atmosphere, the surface of aluminum is prone to forming a thin oxide layer from oxygen in the environment that reacts with the powder during processing; extremely low O2 partial pressure, <10−143 torr, is necessary to prevent oxidation (Kowalski et al., 1992). Increased hydrogen contamination results from hydroxides and stearic acid, which coupled with oxygen contamination can have a detrimental effect on ductility (Lu et al., 2002). A characteristic of cryomilled materials is the
Table 2 Tensile and hardness properties of the Al 5083-plus-B4 C nano-composite plates in the longitudinal orientation. Data on Standard Al 5083 H131 are provided for comparison. Processing
0.2% yield strength (MPa)
Ultimate tensile strength (MPa)
Elongation (%)
Elastic modulus (GPa)
HIP/HSRF HIP/QIF QIF H131a
397 320 318 255
502 471 460 338
0.9 2.0 1.7 14.0
80.5 82 81 71
NC, Nanocrystalline region; CG, coarse-grained region. a Standard Al 5083 H131 referenced from (Horn, 1967).
Hardness (Vickers) NC/CG
NC
150 ± 11 136 ± 9 158 ± 9 –
232 ± 22 214 ± 18 216 ± 16 –
CG 85 ± 7 99 ± 3 100 ± 4 90
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Fig. 4. Grain size distribution histograms for: (a) HIP/HSRF longitudinal axis, (b) HIP/HSRF transverse axis, (c) HIP/QIF longitudinal axis, (d) HIP/QIF transverse axis, (e) QIF longitudinal axis, and (f) QIF transverse axis. Material is Al 5083-plus-B4 C nano-composite.
increased thermal stability due to small dispersoids that result from the presence of the oxygen and nitrogen, giving rise to Zener pinning (Witkin and Lavernia, 2006). The stearic acid used as a process control agent is the primary source of the carbon and hydrogen in the cryomilled-and-blended powder; removal of these is the primary purpose of the degassing step. The hydrogen content decreased ∼330 ppm from the cryomilled-and-blended powder to the final consolidated composite plates; this reduction can be attributed to the degassing step. As a trade-off, however, the degassing also results in grain coarsening of the Al 5083 matrix from ∼25 nm after cryomilling to ∼168 nm after degassing. The mechanical behavior of the composite plates in tension and compression exhibited an asymmetry in yield strength while the Standard H131 plate exhibited some symmetry. The HIP/HSRF plate yielded at a value ∼100 MPa higher in compression than in tension. The QIF plates yielded roughly 60 and 70 MPa higher in compression than in tension for the HIP/QIF and QIF plates, respectively. This asymmetry behavior has been observed before by Ye et al. (2006a,b) in extruded MMC’s and is a generic feature in composites (Clyne
and Withers, 1993). Upon cooling from the processing temperature there is a residual tensile stress field that remains in the composite (Fernández et al., 2004) resulting in a lower tensile yield strength. The compression testing in the forging direction and perpendicular to the forging direction for all three plates show anisotropic behavior that can be explained by the anisotropic microstructure that results during deformation (Skolnik et al., 2008). The fracture surfaces, shown in Fig. 5, exhibit well-defined dimples in the CG bands indicating a ductile fracture mechanism after all three processing methods. In the particulate reinforced NC matrix regions, the fracture morphology is not as distinct as in the unreinforced CG regions with much shallower peaks and valleys. B4 C particles are not apparent at the NC matrix fracture surfaces suggesting that debonding between the matrix and reinforcement is not a predominant fracture mechanism, which is consistent with the behavior observed previously in aluminum reinforced with SiC particulate (Beffort et al., 2007). The HIP/HSRF plate showed higher strength values, with limited ductility, than the plates processed by QIF. The grain size distri-
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bution data shown in Fig. 4 indicate that the distributions in the NC regions are similar in the three plates, and therefore, grain size strengthening does not appear to be the reason for the observed differences in mechanical behavior. Considering the consolidation techniques in more detail, it is important to note that HSRF operates at an extremely high strain rate, 102 s−1 , compared to QIF at 10−2 s−1 . This high strain rate can prevent dynamic recovery or recrystallization during the processing; this is supported by the evidence that the HSRF plates showed a higher aspect ratio (1.9) compared to that of the QIF plates (1.6 and 1.3). The lower aspect ratio in the QIF plates reflects a large portion of recrystallized grains in the QIF plates. As shown in Fig. 4, the grain size distribution
Fig. 6. Work hardening rate for the three plates consolidated using three different methods based on compression results in the transverse and longitudinal orientation. Material is Al 5083-plus-B4 C nano-composite.
widened in the transverse direction and narrowed in the longitudinal direction in the QIF forged plates compared to that in the HSRF plates. To further consider this effect, the work hardening rate can be calculated by =
Fig. 5. Tensile fracture surfaces of (a) HIP/HSRF plate, (b) HIP/QIF plate, and (c) QIF plate. Material is Al 5083-plus-B4 C nano-composite.
d dε
(1)
where is the change in stress divided by the change in strain (Rodriguez et al., 2003), the results of which are shown in Fig. 6 calculated using true stress–strain compression data. The work hardening rate is higher for the QIF plates than for the HSRF plate, which is in agreement with the lower strengths and increased ductility of the QIF plates, suggesting that the QIF plates have lower initial dislocation density. In nanocrystalline materials, recovery or recrystallization may occur via rotation or coalescence of adjacent smaller grains (Zhou et al., 2003). The extent of grain rotation and coalescence is expected to increase during deformation in response to the orientation accommodation between adjacent grains under an applied load (Jin et al., 2004) or driven by stress-assisted grain boundary migration (Jin et al., 2004; Gianola et al., 2006). In a study by Witkin et al. (2005), evaluating microstructural coarsening during thermomechanical processing, they found that recovery or recrystallization occurred quickly after the onset of deformation but was not sensitive to the strain. Therefore, in addition to the strain rate effect, the larger extent of recrystallization in the QIF plates can also be caused by the fact that the QIF plates underwent more deformation steps than the HSRF plates. This view suggests that a single pass high strain rate consolidation method is preferred to achieve high strength while subsequent thermomechanical processing can be used to increase the ductility in the Al 5083-plus-B4 C nano-composite.
R.G. Vogt et al. / Journal of Materials Processing Technology 209 (2009) 5046–5053
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