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Macro- and Microstructural Studies of Laser-Processed WE43 (Mg-Y-Nd) Magnesium Alloy S. SANTHANAKRISHNAN, N. KUMAR, N. DENDGE, D. CHOUDHURI, S. KATAKAM, S. PALANIVEL, H.D. VORA, R. BANERJEE, R.S. MISHRA, and NARENDRA B. DAHOTRE The macro- and microstructural changes in the fusion zone (FZ) of WE43 magnesium alloy processed by a diode-pumped ytterbium (IPG YLS-3000) continuous wave fiber laser for specified processing conditions (from 4.17 to 12.5 9 107 J/m2) were studied. With the aid of computational heat-transfer model, the temperature history and cooling rate were determined for different laser-processing conditions. The computational heat-transfer results of laser-processed samples were used to correlate with microstructures characterized using optical, scanning, and transmission electron microscopies, and electron backscatter and X-ray diffraction analyses. The microhardness measurement was carried out to establish the structure–property relationship, and the results indicated that the minimal hardness variation (1 pct) within laser FZ was due to a constant heat extraction time (0.1 second), narrow variation in grain size (4 to 7 lm), and the type of precipitate (b-phase) formation. DOI: 10.1007/s11663-013-9896-7  The Minerals, Metals & Materials Society and ASM International 2013

I.

INTRODUCTION

DOWNSIZING vehicle weight, fuel economy, overall production cost, product reliability, market sustainability, and service durability are among the key benchmarks of the transportation industry. Toward that end, Mg alloys are widely acknowledged as suitable candidates to serve the purposes of transportation industry. Mg is 36.8 and 78.2 pct lighter per unit volume than aluminum and iron, respectively; and possesses good physical, chemical, and mechanical properties useful for various applications.[1] Various Mg alloys such as AZ (Mg-Al-Zn), AM (Mg-Al-Mn), AS (Mg-Al-Si), AE (Mg-Al-RE), and WE (Mg-Y-RE) have been designed, developed, and successfully used according to specific applications.[1] However, properties such as low fatigue, as well as creep properties, and poor resistance to corrosion and wear have limited their structural applications. The macroscopic mechanical properties of the Mg alloys are the manifestations of the micromechanics of interaction between dislocations and many other microstructural elements present in a metallic material. S. SANTHANAKRISHNAN, formerly with Department of Materials Science and Engineering, University of North Texas, Denton, TX 76203, is now Assistant Professor at the Manufacturing Division, Mechanical Engineering Department, Indian Institute of Technology-Madras, Chennai 600 036, India. N. KUMAR and D. CHOUDHURI, Post-Doctoral Associates, N. DENDGE, S. KATAKAM, S. PALANIVEL, and H.D. VORA, Doctoral Students, R. BANERJEE and R.S. MISHRA, Professors, and NARENDRA B. DAHOTRE, Professor and Chairman, are with the Department of Materials Science and Engineering, University of North Texas. Contact e-mail: [email protected] Manuscript submitted February 26, 2013. Article published online June 11, 2013. 1190—VOLUME 44B, OCTOBER 2013

Precipitation strengthening is a widely recognized means of improving the strength of many commercially available alloys. Temperature-dependent kinetics are crucial in governing the competing aspects (diffusional and thermodynamic) of each element in Mg alloys. In light of this, a precisely controlled, high-power laser can be used to tailor the surface and the corresponding mechanical properties of Mg alloys. Laser processing (LP) generates localized rapid heating followed by rapid cooling concurrently, which in turn results in rapid melting and solidification. In this process, the localized laser–material interaction for short duration (ls to ms) produces conditions far beyond the thermodynamic equilibrium state, thereby resulting in unique surface and mechanical properties. Hence, judicious selection of alloying elements (chemistry and composition), processing technique(s), and processing parameters may be key to optimization of mechanical properties through the introduction of important microstructural components.

II.

BACKGROUND

To enhance the mechanical and surface properties of Mg, elements such as aluminum (Al), Zirconium (Zr), rare earth (RE) elements (scandium (Sc), neodymium (Nd), yttrium (Y), cerium (Ce)), Manganese (Mn), Silicon (Si), and Zinc (Zn) are added. The addition of rare earth elements to Mg suppresses the microporosity formation, which results in enhanced formability.[2] The addition of Y enhances materials strength via solid solution and precipitation strengthening, while the addition of RE elements like Nd promotes primarily precipitation strengthening by various precipitates.[3,4] Thus, such additives impart enhanced strength and METALLURGICAL AND MATERIALS TRANSACTIONS B

favorable creep properties to Mg alloys.[3,4] On the other hand, elements like Zr promote grain refinement, which further contributes to the mechanical strength of the alloy.[3–5] Recently, much attention has been given to studying the rare earth-based WE43 (Mg—4.0 wt pct, Y—3.3 wt pct, and Nd (including Dy, Gd)—0.5 to 0.6 wt pct Zr) Mg alloy because of several advantages it offers over conventional Mg alloys.[2,3,5–8] Precipitation hardening and microstructural refinement are the most preferable mechanisms to enhance the strength and ductility of WE 43 alloy.[8] Aging heat treatment (e.g., T6-type treatment) is carried out to take advantage of precipitation hardening. During such treatments, generally, b¢¢, b¢, and b1 metastable precipitates and equilibrium b precipitate are formed by varying the aging temperature and time. In earlier stage of aging [483 K (210 C)], a metastable, hexagonal close-packed (hcp)-ordered DO19 b¢¢ is formed. Longer aging at 483 K (210 C) is favorable to form an intermediate, body-centered orthorhombic (bco) b¢ (Mg12NdY), precipitate. Equilibrium faced-centered cubic (fcc) b (Mg14Nd2Y) is formed at higher temperatures.[2,6] Various laser sources such as CO2[5] and Nd:YAG[1,9,10] have been used to optimize the processing variables (laser power, scanning speed, shielding gas flow rate, and stand-off distance), to obtain the required mechanical properties of joints in AZ31,[1,5,11] AM60,[5] ZK60,[5] AZ91,[7,9] ZE41A-T5,[1,10,12] and WE43[7,12,13] Mg alloys. Laser processing (LP) results in four distinguishable microstructural regions: fusion zone (FZ), partially melted zone (PMZ), heat-affected zone (HAZ), and base material (BM). Such regions result from the variation in rates of heating, cooling, and solidification characteristics to each of these regions.[1] Furthermore, the majority of the studies[1,9–11] reported that the aspect ratio (penetration depth to width) of the joint (fusion region) and mechanical properties deteriorated with increasing laser scanning speed. With this background, the current study attempted to study the macro- and microstructural changes in the joints of WE43 alloy produced using various LP conditions along with microhardness measurement to establish microstructure–property correlation. A numerical heat-transfer tool ‘‘COMSOL Multiphysics,’’ and experimental characterization tools (optical microscopy (OM), scanning and transmission electron microscopes (SEM, TEM), electron backscattered diffraction (EBSD), X-ray diffraction (XRD), and Vickers microhardness indentor were used to obtain the thermal history, microstructures, and microhardness for different LP conditions, and to establish the corresponding structure–property relationship.

III.

cast, solutionized [798 K (525 C) for 8 hours), and hot rolled to 95 pct reduction in thickness. During the current investigation, LP was performed on the as-received WE43 plates. The plates (600 9 600 9 5 mm3) were cut into 50 9 50 9 5 mm3 size coupons and then polished using a 400 grit SiC paper, followed by cleaning with ethanol in ultrasonic bath before LP. The coupons were subjected to different LP (Figure 1) conditions (Table I). A diode-pumped ytterbium (IPG YLS-3000) continuous wave (CW) fiber laser (1064-nm wavelength) with Gaussian beam of focal spot of 0.6-mm diameter on the surface at a fixed stand-off distance of 504 mm from the surface of sample was used to perform the bead-on-plate experiments. The short transverse cross section for each sample was prepared for microstructural observations and analysis using standard metallurgical techniques. Specimens were etched for about 3 seconds in a mixture of acetic glycol ((20 mL) acetic acid, (1 mL) HNO3, (60 mL) ethylene glycol, and (20 mL) water), and finally washed using ethanol. The optical (Nikon Eclipse ME 600), scanning electron (FEI Qunata 200), and transmission electron (Tecnai TF 20) microscopes were used to reveal the microstructures and distinguish the processed regions from BM. The samples (10 9 5 9 5 mm3) for EBSD were prepared by mechanical polishing followed by electropolishing [40 vol pct H3PO4 and 60 vol pct ethanol at 293 K (20 C)] using voltage of 2 to 3 V for a duration range of 10 to 15 seconds. An FEI Nova NanoSEM 230 (20 kV, 3.1 mA, tilting 70 deg, 1 lm step size) was used to record the EBSD patterns from the fusion/melt region (joint or bead) of LP coupons. The volume fraction of the precipitate within the fusion/ melt region was calculated using ImageJ software. For TEM analysis, circular thin foils were prepared from the BM using conventional dimple grinding and ion-milling procedure. However, owing to limited

EXPERIMENTAL PROCEDURES AND NUMERICAL MODELING

A. Experimental Setup and Procedures WE43 alloy used in the current study was obtained from Magnesium Elektron North America Inc., Madison, IL, in hot-rolled condition. The alloy was METALLURGICAL AND MATERIALS TRANSACTIONS B

Fig. 1—Schematic of laser beam focusing assembly and sample orientation during processing of WE43 alloy. VOLUME 44B, OCTOBER 2013—1191

Table I. Laser Processing Parameters Employed in the Present Experiments

S. No

k

 @T ¼ q_  h1 ðT  T0 Þ  re T4  T40 @z

½2

Laser Power (W)

Scanning Speed (m/s)

Laser Beam Diameter (m)

Laser Energy Density (9 107 J/m2)

q_ ¼

  4gP y0  y exp pd2 2u2

½3

500 1000 1500

0.02 0.02 0.02

0.0006 0.0006 0.0006

4.17 8.33 12.5

k

@T ¼ h2 ð T  T 0 Þ @z

½4

k

@T ¼ h3 ð T  T 0 Þ @n

½5

1 2 3

Table II. Temperature Dependent Density (q), Specific Heat (c), and Thermal Conductivity (k) of WE 43 Alloy[9] T [K (C)]

q (kg/m3)

c (J/kg K)

k (W/m K)

300 (573) 400 (673) 500 (773) 600 (873) 700 (973) 800 (1073) 900 (1173) 1000 (1273)

1840 1785 1730 1675 1625 1575 1525 1410

699 716 734 752 771 790 809 830

51 54 58 62 66 71 76 43

volume of fusion/melt (joint) region, site specific TEM thin foils were extracted from the FZ of LP region subjected to 8.33 9 107 J/m2 laser power with focused ion beam (FIB) in dual beam (FEI Nova 200) SEM. The microhardness measurement was carried out on the transverse cross section of as-received and laserprocessed material. A load of 1.96 N in combination with 10 seconds dwell time was used for microhardness measurement. The sample for hardness measurement was prepared in the same way as for OM. All the measurements were taken at the middle of the laser-processed zone depth. For unprocessed alloy in as-received and aged conditions, the microhardness values represent average of 10 data points. Each indent was at sufficiently large distance to avoid overlapping of the plastically deformed field around each indent. The microhardness values for FZ, PMZ, and HAZ represent an average of two to six data points. Each indent for these three zones was 220 lm apart from the neighboring indent. A Rigaku III Ultima XRD with CuKa of 0.15418-nm wavelength at 40 kV and 44 mA in a 2h range (20 to 90 deg at a step size of 0.02 deg and 2 deg/min) was used to identify and quantify the phases of LP coupons. B. Numerical Modeling A 3-D transient conduction heat-transfer equation (Eq. [1]) with as-used boundary conditions (Eqs. [2] through [5]) and temperature-dependent materials properties (Table II) were used in COMSOL Multiphysics to obtain the temperature history and cooling rate for various LP conditions.       @ ðqcðTÞÞ @ @ ðkðTÞÞ @ @ ðkðTÞÞ @ @ ðkðTÞÞ ¼ þ þ @t @x @x @y @y @z @z

½1

1192—VOLUME 44B, OCTOBER 2013

where q(T), c(T), and k(T) are temperature-dependent density (kg/m3), specific heat capacity (J/kg K), and thermal conductivity (W/m K) of the material, respectively. q_ is the heat flux (W/m2), T is an instantaneous temperature (K), T0 is the room temperature (K), r is Stefan–Boltzmann constant (5.67 9 108 W/m2 K4), e is emissivity (0.75), g is absorption coefficient (0.25), P is laser power (W), d is laser beam diameter (0.6 9 103 m), u is an effective beam diameter (0.4 9 103 m), and h1, h2, and h3 are the heat-transfer convection coefficients at top (100 W/m2 K), bottom (10 W/m2 K), and other sides (20 W/m2 K), respectively. The details of heat-transfer model and corresponding assumptions were reported earlier.[14]

IV.

RESULTS AND DISCUSSION

The optical micrographs (Figures 2(a) through (c); Table III) and previous results[1,5,9,10] indicate the aspect ratio (AR = penetration depth/processed width) of laser-processed zones increases as the laser energy density (LED, J/m2) increases. A linear relationship exists between the laser energy density and aspect ratio (Figure 3 and Eq. [6]), AR ¼ n  LED LED ¼

P d  A v

½6 ½7

where A is the laser irradiated area, v is the scanning speed (m/s), and n is an empirical constant. Key transition zones (Figure 4) such as FZ, PMZ, HAZ, and BM are the consequence of steep temperature variation in various regions (Figure 5) due to the dynamic nature of LP. In all cases, the average peak temperature and cooling rate increase as the laser energy density increases (Table IV). As the laser energy density increases, the shape of laser fusion/melt region changes from semi-circular (4.17 9 107 J/m2) to parabola (8.3 9 107 J/m2) and to hyperbola (12.5 9 107 J/m2), respectively (Figures 2(a) through (c) and 3). Once the energy density increases above 8.3 9 107 J/m2, the mode of transition of the fusion/melt region changed from conduction to keyhole type, causing vaporization and material loss. In fact, with high laser energy density

METALLURGICAL AND MATERIALS TRANSACTIONS B

Fig. 2—Optical micrographs of the trend in transition of FZ in WE43 as function of input laser energy density.

Table III.

Laser Energy Density as Function of Physical Attributes of Heat-Transition Region

Laser Energy Density (9 107 J/m2) 4.17 8.33 12.5

Penetration Depth (mm)

Processed Width (mm)

Aspect Ratio

0.27 0.93 2.08

1.37 2.13 2.96

0.33 ± 0.02 0.76 ± 0.035 1.17 ± 0.05

input and fixed beam focal spot area, a high pressure molten plasma above the FZ is usually created perpendicular to the laser scanning direction. Simultaneously, the material surrounding the fusion/melt region acted as a heat sink and led to rapid self-quenching (>103 K/s) and formation of a keyhole as the footprint of Gaussian energy distribution (Figure 2(c)). However, the keyhole formation, vaporization, and material loss can be minimized by appropriately controlling the LP variables (laser power, scanning speed, beam focal spot size, and so on). Significant variation in microstructure such as equiaxed, columnar, and coarse grains (Figure 4) as a result of nonuniform rates of heating (103 K/s), cooling (104 K/s), and subsequent solidification (103 m/s) can be observed across and along the key transition zones (Table IV; Figure 4). The rapid cooling experienced during LP further leads to significant grain refinement in the FZ. At the fusion boundary, a combination of relatively larger thermal gradient and small growth rate predominantly developed a columnar microstructure. The morphology changes from equiaxed to columnar as the laser energy density decreases (Figure 4). Rapid melting and solidification during LP tend to generate nonequilibrium thermodynamic conditions. This in turn can result in microstructures that are substantially different from those produced by conventional thermal treatment such as the aging process. That the grain size in the FZ varied from 4.0 to 7.0 lm METALLURGICAL AND MATERIALS TRANSACTIONS B

Fig. 3—Aspect ratio as function of input laser energy density indicating the trend in transition from conduction to keyhole mode.

during LP (Table V) is noteworthy. Factors resulting in such grain sizes include, but are not limited to, the presence of Zr particles in WE43, which act as a catalyst for subsequent grain nucleation and the extremely fast cooling rates (104 K/s) of LP, which restrict the tendency for grain coarsening. The presence of Zr-containing particles especially helps in accelerating the grain nucleation rate and inhibits the grain growth process by pinning the grain boundaries. This is evident from the fine grain sizes obtained in the present study after LP. Although the laser energy density varies from 4.17 9 107 to 12.5 9 107 J/m2, overall, only a marginal variation in grain size was observed (Table V). Based on the numerical simulation, the times taken for dissipation of heat (surface temperature/cooling rate) through the linear distance of 20 mm are 100 ms, 60 to 80 ms, and 160 to 250 ms within the surrounding regions of FZ, VOLUME 44B, OCTOBER 2013—1193

Fig. 4—Secondary electron images of as-received and laser-processed WE4 showing grain structure and its variation in various heat-transition regions of laser-processed samples.

Fig. 5—Temperature history of the FZ as function of time in laser-processed WE 43 at various input laser energy densities indicating associated rapid heating and cooling cycles.

PMZ, and HAZ, respectively (Table IV). This suggests an optimal interplay between cooling rate and coarsening kinetics. 1194—VOLUME 44B, OCTOBER 2013

The microhardness data from transition zones (Table V) reveal a decrease in microhardness values compared with the BM (substrate) by ~6 to 15 pct. The METALLURGICAL AND MATERIALS TRANSACTIONS B

METALLURGICAL AND MATERIALS TRANSACTIONS B

VOLUME 44B, OCTOBER 2013—1195

4.17 8.33 12.5

Laser Energy Density (9 107 J/m2)

4.17 8.33 12.5

Laser Energy Density (9 107 J/m2)

4 to 5 5 to 6 6 to 7

Grain Size (lm)

Table V.

9500 16,300 24,500

Cooling Rate (K/s) 905 (1178) 920 (1193) 940 (1213)

Peak Temperature [K (C)] 6350 10,900 16,300

Cooling Rate (K/s)

Partially Melted Zone (PMZ)

670 (943) 720 (993) 780 (1053)

Peak Temperature [K (C)]

6.0 4.0 3.5

Volume Fraction Pct

Fusion Zone

53.4 ± 1.7 57.6 ± 1.0 56.7 ± 1.0

Hardness (kgf/mm2)

12.5 11 10

Grain Size (lm)

51.0 ± 2.2 58.3 ± 0.7 58.4 ± 1.1

Hardness (kgf/mm2)

Partially Melted Zone

18.5 21 22

Grain Size (lm)

56.2 ± 2.8 59.4 ± 2.9 56.4 ± 3.5

Hardness (kgf/mm2)

Heat Affected Zone

35

Grain Size (lm)

63.5 ± 4.4

Hardness (kgf/mm2

As-Received Material

82 ± 5.7

Base Material (Aged) Hardness (kgf/mm2

1900 3300 4900

Cooling Rate (K/s)

Heat Affected Zone (HAZ)

Grain Size, Volume Fraction of Precipitate, and Hardness in Key Transition Zones as Function of Input Laser Energy Density

990 (1263) 1630 (1903) 2450 (2723)

Peak Temperature [K (C)]

Fusion Zone (FZ)

Table IV. Temperature History and Cooling Rates as Function of Input Laser Energy Density Corresponding to Various Heat-Transition Regions

Fig. 6—SEM images of melt region of laser-processed WE43 at various input laser energy densities showing variation in grain size and grain boundary precipitates as function of input laser energy density.

Fig. 7—Microscopy of laser-processed (with 8.33 9 107 J/m3 laser power) and base metals: (a) bright-field TEM obtained close to [1010]Mg zone axis of grain #1 show precipitation at the grain boundaries and within grain (#2) in the laser-processed region, (b) bright-field TEM obtained close to [11 20]Mg zone axis show dislocations in the base metal, (c) SEM image in backscattered mode show intergranular equilibrium-b precipitates and very fine scale (marked with arrows) intragranular precipitates in the base metal, and (d) fine scale precipitates in the base metal are shown with atomic number contrast by high-angle annular dark-fields-scanning transmission electron microscopy (HAADF-STEM).

flow stress and mechanical strengths are governed by a set of interacting variables. These variables depend on the efficacy of solutes in terms of solid–solution strengthening, dislocations, second-phase particles, and 1196—VOLUME 44B, OCTOBER 2013

geometrical constraints posed by grain boundaries that obstruct the dislocation motion. Apparently, irrespective of the refinement of grains in laser-processed samples compared with as-received material, the METALLURGICAL AND MATERIALS TRANSACTIONS B

hardness values of laser-processed samples have dropped considerably (~6 to 15 pct). This indicates a weaker dependence of flow stress on the grain size.[15–20] In light of this, the strengthening levels appeared to be extremely sensitive to precipitation. Moreover, previous extensive studies on Mg-Y-Nd system[15,21–23] suggest that precipitation is very important in determining the final hardness values. In light of this, grain size appeared to have less impact on strengthening the laser-processed WE43 alloy. Evidently, the volume fraction of precipitate decreased as the laser energy density increased (Figure 6; Table V). The level of hardness values indicated that these precipitates possibly belong to the b phase. Based on the temperature history experienced during LP (Figure 5) and secondary electron images of the FZs (Figures 6(a) through (c)), the FZ microstructure appeared to resemble with that of an as-cast microstructure (i.e., rapid melting followed by rapid cooling). Such microstructural features in the fusion/melt region of laser-processed material were further probed with TEM and compared with the microstructure of as-received sample. Bright-field TEM images displayed in Figure 7 provide a comparison between the microstructure in the FZ of a region processed with 8.33 9 107 J/m2 (Figure 7(a)) and the as-received material (Figure 7(b)). Vertical striations noted in the bright-field image of FZ (Figure 7(a)) are an artifact of focused-ion-beam milling. Nonetheless, this bright-field image of the laser-processed region depicts precipitates, in darker diffraction contrast at the grain boundary and some within grains. Figure 7(a) shows the morphology of intragranular precipitates, which are either cuboidal/rectangular (Grain 1) or globular (Grain 2). The grain boundary precipitates are also indicated with an arrow in Figure 7(a). Such precipitate morphologies are consistent with TEM observations on as-cast WE43 by Rzychon and Kielbus.[24] These authors also indicated that the grain boundary precipitate is b (Mg14Nd2Y), while the intragranular precipitates have Mg41Nd5, or Mg24Y5 compositions. The details regarding the crystal structure, lattice parameters, and composition of the precipitates observed in the laser-processed samples in the present case are the subject of ongoing investigation. Furthermore, the authors emphasize that the focus of the current study is to establish microstructure–mechanical response correlations at a macroscopic level. The distribution and morphology of precipitates in FZ and as-received sample are also compared (Figure 7). Figure 7(b) is a bright-field TEM image showing the dislocation substructure within a grain of the as-received WE43 sample. A backscattered SEM image from the same sample, exhibiting strong atomic mass contrast (Z-contrast), is shown in Figure 7(c), with large b-type grain boundary precipitates (marked by the arrow in Figure 7(c)) and finer scale intragranular precipitates. Figure 7(c) shows very fine scale intragranular precipitates within the top grain, and precipitation on vein-like structures in the lower grain. The fine vein-like structures were also seen while using HAADF-STEM imaging (Figure 7(d)). The HAADF-STEM images exhibit the local enrichment of heavy elements (e.g., Nd, Y) in brighter contrast, thus highlighting the precipitates. METALLURGICAL AND MATERIALS TRANSACTIONS B

Compared with the hot-rolled microstructure, FZ of laser-processed sample shows discrete precipitation within the grains. Large intragranular cuboidal/rectangular and globular precipitates are observed (Figure 7(a)). The above differences in microstructural features between the laser fusion/melt region and as-received material may influence their respective hardness values. Dislocation motion in the as-received material can be retarded by the preexisting dislocation substructures (Figure 7(b)) and intragranular precipitates (Figure 7(c)). On the other hand, larger intragranular precipitates separated by several hundred nanometers may not serve as effective barriers to dislocation propagation. Thus, comparison of precipitate distribution and morphology in FZ and as-received sample suggests that the microstructure in laser-processed sample will result in lower strength than that of the as-received material itself. This analysis is also consistent with the measured hardness values of the FZ and as-received BM. Irrespective of the presence of large volume fraction of nonstrengthening b-phase, in general, the impact on hardness of the laser fusion/melt region is marginal.[2,21] At higher laser energy density, the interparticle distance is greater, thereby increasing the inter-obstacle spacing to dislocation motion. In addition, the cooling curves (Figure 5) indicate that greater time is required to reach room temperature as the laser energy density increases. This has two implications. First, it may lead to greater elemental diffusion and higher coarsening tendencies for higher laser energy density specimens, which in turn may have lower hardness values. Apart from the precipitate size and shape, the cooling rate would also have considerable impact on the precipitate habit plane that controls the strength values.[22,25] Second, the cooling rate does not appear to be fast enough to retain a high amount of solute in solid solution. The effect of laser energy density on the evolution of phases is also studied by semi-quantitative XRD analysis. A semi-quantitative phase analysis was conducted from XRD patterns by measuring the sum of intensities of the peaks for the respective phases and normalizing it with the sum of all the peaks. The results indicate that the phase fraction of the rare-earth oxides (yttrium oxide) increases with increase in laser energy density from 11.6 pct for the lowest laser energy density to 13.3 pct for the highest laser energy density (Figure 8), which in turn depletes the solutes influencing the hardening ability and thereby reducing the strength. In all cases, the presence of Mg is confirmed. The MgOphase fraction increases as the laser energy density increases from 8.3 pct for lowest laser energy density to 9.3 pct for the highest laser energy density. Moreover, at higher laser energy density when the laser beam moves to the next location, more heat needs to be conducted/ dissipated from the previous location, thereby raising the temperature and accelerating the coarsening kinetics. The sustenance of faster cooling rate until lower critical temperature could induce higher defect diffusivity and higher precipitate concentration along the grain boundaries, thereby decreasing hardness values. Furthermore, during LP, elemental segregation can occur VOLUME 44B, OCTOBER 2013—1197

within the interdendritic regions and decrease in hardness values.[24] Hence, generating the right type and characteristics of precipitate to achieve higher strength is important. Mg has hcp crystal structure that possesses limited slip systems such as (0001) h11 20i basal slip, (10 12) h1011i pyramidal twinning, and more complex slip on prismatic (10 10) and pyramidal (10 11) planes.[26,27] The lowest

Fig. 8—XRD patterns of laser-processed WE43 at various input laser energy densities.

critical resolved shear stress (CRSS) is required for basal slip as compared with other slip systems. The CRSS of twinning and prismatic slip are larger compared with CRSS of basal slip.[28] The yield strength and deformation of Mg depend greatly on type of texture. The deformation of basal plane is more favored, if texture of basal plane is aligned to 45 deg. When basal plane texture is random, more grains are oriented at an angle 45 deg. This indicates high deformation capacity in the case of random texture. Other than precipitate details (phase, morphology, distribution, etc.) and grain size, texture-induced anisotropy may also influence mechanical strength. Thus, texture-related influences were examined with EBSD observations. Figures 9 and 10 present orientation image maps (OIMs) and orientation distribution of {0002}, {1010} and {1120}, respectively. These figures compare EBSD results of as-received WE43 (Figures 9(a) and 10(a)) with laser-processed samples at laser energy densities of 4.17 9 107 J/m2 (Figures 9(b) and 10(b)), 8.33 9 107 J/m2 (Figures 9(c) and 10(c)), and 12.5 9 107 J/m2 (Figures 9(d) and 10(d)). To facilitate one-on-one comparisons, EBSD scans and hardness measurements were performed on the same cross section of each sample. Furthermore, EBSD results of only fusion/melt zone and not PMZ and HAZ of the laser-processed region were presented because the FZ

Fig. 9—OIM comparing grain sizes and grain orientations of (a) base WE43 as-received and FZs of laser-processed samples with energy densities, (b) 4.17 9 107 J/m2, (c) 8.33 9 107 J/m2, (d) 12.5 9 107 J/m2, (e) color-coded stereographic, and (f) reference frame attached to the plane of the OIMs, indicating transverse (TD) and rolling (RD) (Color figure online). 1198—VOLUME 44B, OCTOBER 2013

METALLURGICAL AND MATERIALS TRANSACTIONS B

 and {1120}  orientation distributions showing Fig. 10—{0001}, {1010} the effect of laser processing on texture: (a) base WE43 as-received and FZs of laser-processed samples with energy densities, (b) 4.17 9 107 J/m2, (c) 8.33 9 107 J/m2 and (d) 12.5 9 107 J/m2. Color-coded intensity scale for each processing condition is also shown (Color figure online).

forms the bulk/majority of the processed region, and therefore may have significant influence on mechanical strength. OIMs presented in Figure 9 clearly show that grain size decreased in laser-processed samples, which is consistent with the SEM results (Figures 4 and 6). Importantly, the OIMs also revealed that LP caused significant differences in grain orientation distribution compared with as-received WE43. Figure 9(a) indicates that grains in as-received material are predominantly oriented close to {10 10} and {11 20} poles (by comparing OIM with color-coded inverse pole figure triangle in Figure 9(e)). Such preferred grain orientation is related to basal texture typically displayed by Mg-based alloys. Corresponding pole figures of the as-received material (Figure 10(a)) indicate a near-basal texture. OIMs of the FZs in laser-processed regions (Figures 9(b) and (c)) compared with as-received BM show significantly larger populations of grains oriented close to {2 poles in addition to {10 10} and {11 20}. The latter observation indicates that the as-solidified microstructures (after LP) have reduced texture strength. Since OIMs merely provide a qualitative depiction of texture, a quantitative description was obtained by plotting orientation distribution of {0002}, {1010} and {11 20}) in pole figures (Figure 10). Orientation distributions were obtained via standard harmonic series expansion available in TSL software. Color map METALLURGICAL AND MATERIALS TRANSACTIONS B

shown next to the pole figures depict relative density of poles in the TD–RD frame of reference, and the numbers corresponding to each color indicate texture strength (e.g., red signifies maximum texture strength). A few key observations were noted from the pole figures presented in Figure 10. The maximum texture strength of as-received material (Figure 10(a)) is 38.2, while those of the laser-processed samples at 4.17 9 107 J/m2 (Figure 10(b)), 8.33 9 107 J/m2 (Figure 10(c)), and 12.5 9 107 J/m2 (Figure 10(d)) are 5.8, 3.7, and 3.1, respectively. The maximum texture strength of laserprocessed samples decreased with increasing laser energy density, suggesting that preferential grain orientation is considerably reduced after LP in comparison with as-received material. Comparison of the distributions of {0002}, {1010}, and {1120} poles of the as-received WE43 (Figure 10(a)) and the laser-processed samples (Figures 10(b) through (d)) further confirmed that preferential grain orientation is significantly reduced after LP in comparison with as-received WE43. Such distributions in the laser-processed samples become increasingly random with an increase in laser energy densities. The random distribution of grains in FZ (Figures 9 and 10) greatly increases the likelihood that several grains (or the Mg HCP lattices) will be favorably oriented with respect to the loading direction (is normal to the pole figures in TD–RD frame of reference) so as to permit dislocation mobility. Therefore, by experimentally determining the angle (F) between the normal to a plane and loading directions, and knowing that CRSS (for dislocation slip) is nonzero for 0 < F < 90, one can reasonably assess if a grain will deform plastically.[29] Using the {0001} orientation distributions for laser-processed samples (Figures 10(b) through (d)), it was found that the majority of the {0002} basal planes were oriented such that 30 < F < 70—a substantial fraction of grains may experience dislocation activity in basal planes. Noting that dislocation slip in basal plane (e.g., {0001} h1120i slip system) in Mg-based alloys has a significant contribution to plastic deformation of such alloys, several grains in FZ likely will be amenable to plastic deformation. Thus, the presence of favorably oriented grains, minimal barrier to dislocation activity by large sparsely distributed precipitates, and weak grain size effect severely reduce the strength of laserprocessed region.

V.

CONCLUSIONS

As the laser energy density increased (4.17 to 12.5 9 107 J/m2), both the penetration depth (0.2 to 2.0 mm) and aspect ratio (0.3 to 1.2) increased, which in turn changed the mode of material processing from conduction to keyhole. A larger thermal gradient and smaller growth rate were observed in the FZ that significantly developed a columnar microstructure. The transition zones (fusion, partially melted, heat-affected) in laser-processed samples reveal a decrease in hardness (~4 to 12 HV) close to that of the BM, but substantially less than the as-received material [cast and solutionized VOLUME 44B, OCTOBER 2013—1199

at 798 K (525 C) for 8 hours, and then hot rolled]. From TEM observations, the grain boundary precipitates were identified as b-type, which influenced the hardness values. The reduction in hardness after LP compared with as-received material was significantly influenced by the lesser variation of heat extraction time (20 to 100 ms), rapid cooling rate (~104 K/s), minimal grain size variation (4 to 7 lm), b-phase formation, and a weak texture (30 to 90 deg).

ACKNOWLEDGMENTS Authors express their sincere thanks to Dr. Junyeon Hwang for his help in getting the as-received WE43 base matrix TEM image and the Center for Advanced Research Technology (CART) for providing the SEM, XRD, and TEM facilities. The authors are grateful to Dr. Bruce Davis of Magnesium Elektron NA for providing the WE43 alloy.

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METALLURGICAL AND MATERIALS TRANSACTIONS B