Porosity characterization of additively manufactured transparent MgAl2O4 spinel by laser direct deposition

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Porosity characterization of additively manufactured transparent MgAl2O4 spinel by laser direct deposition John M. Pappas, Xiangyang Dong∗ Mechanical and Aerospace Engineering, Missouri University of Science and Technology, Rolla, MO, 65409, USA

A R T I C LE I N FO

A B S T R A C T

Keywords: Transparent ceramics Magnesium aluminate spinel Laser direct deposition Additive manufacturing

Polycrystalline magnesium aluminate (MgAl2O4) spinel has established itself as one of the leading candidates for use as transparent, high strength ceramics. In this study, additive manufacturing of spinel ceramics by laser direct deposition was investigated with a particular focus on porosity characterization. Residual porosity was the most significant factor hindering transparency of spinel ceramics. Hence, this paper systematically studied how processing conditions and initial powder particle size affected porosity and densification of laser direct deposited spinel samples. Scan speed, laser power, and powder flow rate, along with powder size, were all shown to have substantial influences on density and porosity of produced parts. High density bulk spinel ceramics, with average densities of nearly 98%, were produced at a low powder flow rate of 0.58 g/min. In addition, the pore size distribution was studied under various processing conditions as it was closely related to transmission properties of transparent spinel at different wavelengths. The reduction in average porosity was mainly attributed to a significant reduction in large pores (> 30 μm) whereas less obvious effects were found on smaller pores. An increase of scan speed resulted in an initial increase in density, followed by a subsequent decrease at scan speeds higher than 2000 mm/min. No obvious difference in porosity was observed between pulsed mode and continuous mode during laser deposition. The use of nanoparticles in initial MgO–Al2O3 mixtures for deposition also promoted densification of the deposited samples. A systematic understanding of residual porosity in this study can help process optimization to minimize porosity in additively manufactured transparent spinel ceramics, potentially eliminating the needs of time-consuming and expensive post-processing procedures.

1. Introduction Transparent ceramics are of great interest for applications such as scratch and chemical resistant optical elements, and transparent armor. High hardness, among other excellent mechanical properties allow for thinner and lighter armor or spacecraft windows than now common multi-layer laminated glasses, which are especially useful for applications in vehicles or aircraft [1]. Transparent polycrystalline materials, including magnesium aluminate (MgAl2O4) spinel, have significant advantages over single crystal materials (e.g. sapphire) due to the potentials for direct production of near net shape components [2]. Net shaped fabrication of transparent ceramics significantly reduces time consuming and expensive post processing requirements. Additive manufacturing (AM) of transparent components can be an effective way to produce near net shaped ceramic parts with minimized requirements for polishing [3], while fabricating complex geometries that would be difficult to manufacture with traditional methods. Compared with cubic polycrystalline ceramics (e.g., spinel in this

∗

study), non-cubic transparent ceramics (e.g., Al2O3 ceramics) present great challenges for producing quality parts with high transparency [4]. Due to birefringence scattering, non-cubic polycrystalline materials should have a grain size of less than one tenth of the optical wavelength to obtain high transparency [5]. However, controlling grain refinement is difficult in ceramic AM processes [6], especially to the point of submicron grain sizes required for optical transparency. Grain refinement in additively manufactured non-transparent ceramics can be achieved via the addition of dopants [7]. However, dopant type and amounts would have to be very carefully chosen to avoid optical losses due to secondary phase formation with different refractive index [8]. Spinel and other cubic polycrystalline ceramics show a primary advantage among other transparent ceramics as they have no inherent birefringence. Owing to a cubic crystalline structure, spinel ceramics are optically isotropic and have high transmission values through a wide range of grain sizes [9]. However, it should be noted that with even the cubic polycrystalline ceramics, the resultant microstructure of fabricated transparent ceramics is still very important and can greatly

Corresponding author. E-mail address: [email protected] (X. Dong).

https://doi.org/10.1016/j.ceramint.2019.11.164 Received 17 September 2019; Received in revised form 12 November 2019; Accepted 19 November 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

Please cite this article as: John M. Pappas and Xiangyang Dong, Ceramics International, https://doi.org/10.1016/j.ceramint.2019.11.164

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Fig. 1. Illustration of residual pores limiting transparency of printed spinel ceramics showing: (A) An opaque sample over a grid of Missouri S&T logos, (B) Porosity of a typical region of the sample, (C) A translucent sample.

porosity started to increase. The initial decrease was attributed to an increase in melt depth and duration as well as acoustic flow aiding the natural buoyancy of gas bubbles. Similarly, Ning et al. [18] showed that porosity of Inconel 718 structures produced by laser direct deposition could be significantly reduced through addition of ultrasonic vibration. The primary mechanisms for reduction in porosity were reported as ultrasonic vibration induced acoustic streaming and cavitation. Although there were extensive studies on the effects of processing parameters on porosity during AM of metals and non-transparent ceramics, their effects during AM of transparent ceramics are not well understood, particularly for laser direct deposition. The only identified work on AM of transparent ceramics was direct ink writing (DIW) of YAG transparent ceramics [19]. DIW, as an indirect AM approach, required additional sintering-based procedures to produce final parts. Thus, the pore formation was governed by diffusion-based densification process, which was relatively well understood from traditional sintering of transparent ceramics [20–22]. In contrast, this paper, based on a new single step AM process by laser direct deposition, is the first research work to study its effects on porosity within the obtained fully meltgrown transparent polycrystalline spinel ceramics. The pore formation must be well understood to produce highly transparent spinel ceramics. Thus, the focus of this study is to determine how processing conditions including laser power, laser operation modes, scan speed, powder flow rate and spot size affect residual porosity in the laser direct deposited bulk spinel ceramics. In addition to the average porosity and density, the pore size distribution is also characterized as the transmission of the printed spinel samples [23] also depends on pore size. This study also investigated the effects of initial powder sizes on the fabricated spinel samples, where two MgO powder precursors with different particle sizes were investigated. The range of processing parameters were determined by a single-track parametric study conducted to obtain continuous deposition of bulk parts.

affect both transparency and mechanical properties of produced parts. For transparent ceramics, the critical microstructural aspects, especially porosity and microcracking that negatively affecting the part transmittance, need to be considered. In the meantime, factors including microcracking, grain size, shape and orientation also play a vital role in determining the mechanical properties of the fabricated parts in structural or armor applications. This study will focus on the porosity, one of the most important microstructural factors [10] affecting both transparency and mechanical properties of laser direct deposited spinel ceramics. Residual pores form a secondary phase with significantly different refractive index from the primary spinel phase [11]. Light passing through a ceramic with constantly changing refractive index results in significant light scattering, rendering the component either translucent or opaque depending on the thickness of the part and prevalence and size of pores. To obtain highly transparent ceramics, porosity should be less than 0.1% with pore sizes less than 100 nm [12], and less than 0.01% porosity is required for glass-like transparency [13]. Therefore, residual porosity must be minimized as much as possible. The negative effects of porosity on transparency are clearly illustrated in Fig. 1. Fig. 1(A) shows a laser direct deposited spinel sample that was mirror polished on both sides to a thickness of 0.7 mm. However, the sample was opaque to visible light, making it difficult to see the Missouri S&T logo placed directly behind. The observed opacity was a direct result of relatively high porosity within the sample as shown Fig. 1(B). At a lower porosity level, a translucent spinel part can be obtained as shown in Fig. 1(C). This clearly illustrates why it is imperative to study how processing conditions will affect porosity, and hence transparency, of laser direct deposited spinel ceramics. Previous studies, focusing on additively manufactured metal and non-transparent ceramic parts, investigated residual porosity to increase densification and improve mechanical properties. Kuriya et al. [14] studied the effects of laser power on porosity for directed energy deposited Inconel 718. It was shown that increased laser power resulted in a significant reduction in gas porosity, which was attributed to longer solidification time allowing large gas bubbles to discharge out of the melt by buoyant forces. Mahamood et al. [15] studied the effects of laser power and scan speed on laser deposited titanium alloy. It was found that increasing laser power and decreasing scan speed resulted in reductions in total porosity. However, at high laser power and high scan speed, the average pore size increased drastically, which was attributed to the increased laser power reducing the number of unmelted particles while the high scan speed limited entrapped gas removal due to rapid solidification. Susan et al. [16] found that a larger porosity content in the powder particles led to a higher porosity in laser direct deposited stainless steel parts. Laser remelting (without powder flow) after deposition significantly reduced porosity content of the deposition. Yan et al. [17] studied the effect of ultrasonic vibration during laser direct deposition of alumina-zirconia eutectics. It was shown that increasing vibration power initially reduced porosity, but after a certain point

2. Experimental procedure 2.1. Materials The materials used in the bulk part deposition include two types of MgO powders from Sigma-Aldrich with a purity greater than 99%, and average particle sizes less than 44 μm and 50 nm, respectively. Each type of MgO powders was mixed with Al2O3 powder from Almatis with a purity greater than 99.5% and D-50 particle size of about 2.2 μm. CoorsTek AD-998 alumina substrates with 99.8% purity were used considering their thermal expansion compatibility with deposited spinel. Powders were mixed to the spinel stoichiometric ratio in a ball mill for 2 h with acetone solvent and alumina grinding media to break up powder agglomerates. Following removal of grinding media, acetone solvent was removed in a Buchi R124 rotary evaporator. To ensure that the mixture was free of organic contaminants and residual moisture, the powder was calcined at 600 °C for 8 h in a Lindberg 2

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Fig. 2. SEM micrographs illustrating powder morphology of mixtures prepared with MgO powder having an as-received average particle size of: (A) Less than 44 μm and (B) Less than 50 nm.

motorized table was used for three-axis positioning and movement of the substrate. A Powder Motions Lab X2W powder feeder was used for powder delivery [24]. High purity argon gas was used to convey the prepared powder mixture. Powder was guided to the melt pool through an alumina tube with inner diameter of 1.6 mm. A one-tube powder delivery setup was used, with powder flow parallel to the scan direction, and at an angle of approximately 70° from the substrate. Fig. 4(A) shows a schematic diagram of the fabricated thin wall structures and approximate dimensions. Thin-wall structures were built using a reciprocal scanning motion shown in Fig. 4(B). After each track, the laser head would increment in the z-direction by Δz, which was equivalent to the layer thickness. Laser spot size was maintained at 2.5 mm diameter throughout this study. To determine effects of processing conditions on porosity of bulk spinel ceramics, parametric studies were performed by varying scan speed, laser power, laser operation modes, powder flow rate, and powder spot size as summarized in Table 1. The powder spot size, maintained at 5 mm except for the parametric study of its effects on porosity, was represented by the crosssectional diameter of delivered powder at the laser focus spot. The printed thin wall specimens were sectioned on a low speed diamond saw and mounted in an acrylic matrix for polishing and microstructural characterization. Cross-sectional micrographs and measurements were taken on a Hirox KH-8700 digital microscope.

furnace. It is worth noting that in this study, bulk single-phase spinel ceramics were directly synthesized from the prepared Al2O3–MgO mixtures during laser direct deposition via melt-grown methods. Fig. 2(A) and Fig. 2(B) show the powder morphology of the Al2O3–MgO mixtures prepared with the initial average MgO particle sizes less than 44 μm and 50 nm, respectively. From Fig. 2(A) it is clearly seen that the particle size of MgO was significantly reduced by the ball milling procedure. On the other hand, Fig. 2(B) shows that the nano-MgO particles tended to clump together to form relatively small agglomerates. In the meantime, the average particle size of alumina in neither of the prepared mixture was significantly affected by the ball milling procedure. The powder used in the single-track study was commercially available spinel nano-powder (Sigma-Aldrich) with particle sizes less than 50 nm. A powder bed was prepared on the substrate with a thickness of 0.4 mm. The spinel powder was mixed with high purity acetone to create a suspension and then was evenly applied onto the substrate. Acetone was then removed by heating to 120 °C before the parametric single-track study.

2.2. Experimental setup and characterization The experimental setup of laser direct deposition is shown in Fig. 3, including a Convergent Energy Arrow Ultimate 1.7 kW CO2 laser operating at 10.6 μm wavelength. A computer numerically controlled

3. Results and discussion 3.1. Characterization of single-track geometries The single-track studies were used to guide the selection of proper processing conditions to achieve efficient material deposition for further printing thin wall specimens. Our preliminary studies showed significant difficulties, mainly affected by the laser scan speeds, in printing thin wall structures of sufficient quality for further analysis of porosity and density due to balling or poor print resolution. Thus, the single-track study was carried out first to guide the selection of processing parameter range, in particular, scan speeds, that provided acceptable deposition quality prior to printing bulk structures. Once found, this range of parameters was used to print bulk structures for determination of the effect of processing conditions on residual porosity and densification. Single-track beads were deposited via varying scan speed across a powder bed. For this study, the laser was operated in continuous wave (CW) mode with a laser spot size of 2 mm. Processing parameters are summarized in Table 2. Single-track geometries were characterized by bead width (w), depth (d), height (h), and contact angle (θ) as shown in Fig. 5. The depth was measured from the top of the deposition to the bottom of the melt pool, whereas the height was from the substrate to the top of deposited bead. Contact angles were measured on both sides of the sample, and their average values were then calculated. Fig. 6 shows the effects of scan speed on the cross-sectional

Fig. 3. A schematic diagram of the experimental setup for laser direct deposition. 3

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Fig. 4. Schematics of printing processes: (A) Fabricated thin wall structures, and (B) Build pattern for the thin wall structures.

significantly increased surface roughness and would necessitate additional post processing to obtain a useable part. This is especially true if optical transparency is desired. Meanwhile, by using a powder bed layer thickness of 0.4 mm (a middle range value for the layer thicknesses used in bulk deposition process), the results were sufficient to guide the selection of the processing range for laser direct deposition. Based on the single-track studies, it was possible to select the proper range of scan speeds that allowed for quality deposition of thin wall structures and subsequent analysis of residual porosity.

Table 1 Processing parameters used for fabrication of bulk spinel ceramics. Scan Speed (mm/min)

1000, 2000, 3000, 4000, 5000

Laser Power (W) Operation modes Powder Flow Rate (g/min) Powder spot size (mm)

275, 485 Continuous wave, pulsed 0.6, 0.9, 1.2, 1.6, 2.1 5 mm, 1 mm

Table 2 Processing parameters for single-track deposition of spinel ceramics. Scan Speed (mm/min)

750–15000 mm/min

Laser Power (W) Laser Spot Size (mm) Powder Bed Thickness (mm)

275 2 0.4

3.2. Porosity and density characterization of spinel samples 3.2.1. Effects of laser scan speed The effect of scan speed on porosity and density of fabricated parts was first investigated. The scan speed was varied from 1000 mm/min to 5000 mm/min, at a constant laser power and powder flow rate of 275 W and 2.1 g/min, respectively. It is worth noting that laser energy density (LED) was maintained at a constant level. As each layer deposited had a finite thickness, LED was represented by volumetric energy density E (J/mm3), calculated as the laser energy per volume material deposited [27] through E = P /(vwt ) , where P is laser power (W), v is scan speed (mm/s), w is laser spot size (mm), and t is layer thickness (mm). Increasing the scan speed would lower the layer thickness accordingly in this study. At a scan speed of 1000 mm/min, the layer thickness was 0.8 mm. After increasing the scan speed to 2000 mm/min, the layer thickness was reduced to 0.4 mm. The calculated volumetric energy density was maintained constantly at 82.5 J/ mm3 for each scan speed tested. Thus, the volumetric energy density would not affect the porosity within the printed samples. Porosity analysis was performed on the cross-section of the sample. The sample was sectioned at a plane through the thickness of the sample, along the build direction, and at the middle of the sample along the scan direction. Due to the variations in processing conditions for the

geometries of deposited tracks. All four variables generally displayed a decaying logarithmic scan speed relationship until the point where balling started to occur. The single-track study showed that single-track beads printed within the range of 1000 mm/min to 5000 mm/min resulted in a continuous, high quality track as shown in Fig. 5(A). At scan speeds higher than 5000 mm/min, single-track beads were very inconsistent with considerable balling observed as shown in Fig. 5(B). Balling occurred when the melt failed to wet the substrate and surface tension pulled the melt into a spherical shape [25,26]. Due to the effect of surface tension, the produced balls showed an increased height to width aspect ratio compared to continuous tracks. As a result, larger fluctuations in both bead height (Fig. 6(C)) and contact angle (Fig. 6(D)) were observed. On the other hand, scan speeds below 1000 mm/min resulted in a continuous track, but preliminary tests showed that print resolution was poor. Poor print resolution

Fig. 5. The optical micrographs of deposited single-track beads showing the measured melt pool geometries: (A) shows a continuous track at a scan speed of 2300 mm/min; (B) shows the balling observed at a scan speed of 6000 mm/min. 4

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Fig. 6. Effects of scan speed on single-track bead geometries: (A) Width, (B) Depth, (C) Height, and (D) Contact angle.

Fig. 7. Densification characterization of the printed samples: (A) shows a typical cross-section with typical defects observed; (B) and (C) summarize the measured relative density and porosity, respectively. (D) shows the post-processed samples printed at scan speeds from 1000 mm/min to 5000 mm/min.

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exhibited large pores and shrinkage cavities, which were not as prevalent in lower layers in Fig. 8(A). The primary reason for lower density in this region was due to the large shrinkage cavities and lack of fusion pores developed near the end of the deposition, identified as large, irregularly shaped cavities produced [30]. As no new layers were deposited above the top regions, the high porosity observed within this region was caused by insufficient liquid phase replenishment during solidification [31]. During cooling of the interior portion of the melt, the material attempted to shrink but was constrained by the outer surface. This resulted in an inability of the liquid phase to completely fill the melt pool [31] and resulted in the shrinkage cavities. Meanwhile, remelting of previous deposited layers would also contribute to reducing porosity, which was related to the total laser heat input of each layer during deposition. As no remelting would occur for the top regions, a lower total laser heat input would be expected compared to that for previously deposited layers. This was further evidenced by the porosity observed before and after further laser heat treatment on the top regions of the printed samples. With no powder feeding, ten additional laser scanning passes were applied on the samples printed at a scan speed of 1000 mm/min in Fig. 8(A). This would allow remelting of top regions. It was seen in Fig. 8(B) that due to the introduced laser heat treatment, the porosity near the top regions was greatly reduced, even lower than that within the lower regions. This suggested the benefits of more laser heat input to further densify the printed samples.

top and bottom layers of thin wall specimens, the middle area of the cross-section was analyzed to obtain the typical microstructural characteristics. Fig. 7(A) shows the area used for the porosity characterization on a typical sample and highlights common defects observed in the printed spinel samples. The image in Fig. 7(A) was stitched together from several optical micrographs using an image stitching plugin in FIJI image analysis software [28]. Laser direct deposited spinel bulk samples typically exhibited a relatively uniform porosity level within the central region as seen in Fig. 7(A), indicating a consistent deposition process. On the other hand, the regions near the part edges showed significant variation in the size and frequency of pores or shrinkage cavities, the formation of which was governed by different mechanisms. These were especially noticeable for the top and bottom regions of the sample in Fig. 7(A). In the top region, it was mainly caused by rapid cooling of the melt when laser energy input ceased, as discussed in more detail below, and thus showed more shrinkage cavities. The bottom region was mainly influenced by interactions with different composition in substrates. Thus, to characterize the typical deposition process, the middle cross-sectional area of samples was analyzed to determine the porosity level. Meanwhile, to measure the corresponding density, the edges of samples were first ground down to obtain cubeshape samples in Fig. 7(D). Density of the post processed samples were then measured using Archimedes’ Principle, with the results summarized in Fig. 7(B). The relative density first increased and then decreased at increasing scan speeds. Initially, large layer thickness produced at relatively low scan speeds of 1000 mm/min limited gas removal from the melt due to long distances that gases had to travel to escape. As scan speed increased to 2000 mm/min, the layer thickness and melt depth decreased accordingly, resulting in shorter distances for pores to the surface of the melt at still high melt pool temperatures. Hence, a higher densification level was observed at 2000 mm/min. At higher scan speeds, layer thickness continued to decrease, which developed a much shallower melt pool. The increase in scan speed resulted in significantly faster cooling rates as was observed in previous studies [29]. At scan speeds in excess of 2000 mm/min melt viscosity increased much more quickly and reduced the time that gasses had to escape by buoyant force. The increase in viscosity and solidification rate dominated the gas-melt interaction, resulting in increasingly lower densities with increased scan speed over 2000 mm/min. It is worth noting that the upward trend in the relative density after 4000 mm/min was attributed to the increased connectivity in pores. Compared to the samples printed at 4000 mm/ min, the higher porosity for the samples printed at 5000 mm/min led to an increased connectivity in the residual pores. As the relative density was measured using Archimedes’ technique, the pores exposed to the sample surface cannot be accurately captured in the density measurement. With increased pore connectivity, more pores would be revealed on the surface of the post-processed samples in Fig. 7(D), yielding an artificially high-density value seen in Fig. 7(B). For this reason, the actual density at the scan speed of 5000 mm/min was expected to be lower than the measured value. To more accurately capture the porosity distribution, the densification analysis was complemented by direct measurements of porosity on the cross-section of the printed samples. Porosity results in Fig. 7(C) further explained the density change observed above. Increasing scan speed from 1000 to 2000 mm/min resulted in a slight decrease in porosity, which corresponded to the increase in density shown in Fig. 7(B). At scan speeds higher than 2000 mm/min, porosity gradually increased until reaching a maximum 14.5% at 5000 mm/min. It is worth noting that without post-processing introduced above, contradicting results were indeed observed in the measured porosity from the middle section and density for the whole samples, i.e., simultaneously increased porosity and density observed at higher scan speeds. This was mainly attributed to difference in porosity formation mechanisms between the middle section (as explained above) and the top regions as seen in Fig. 8. The top layers of printed spinel ceramics

3.2.2. Effects of laser power on residual porosity Laser power was one of the most influential processing parameters on both deposition quality and residual porosity. It was found that with increased laser power and more heat input, significant reductions in both pore size and frequency were made possible. Fig. 9(A) and Fig. 9(B) give a direct comparison of spinel samples printed at laser powers of 275 W and 485 W, respectively. To allow direct comparison of the samples printed at different laser power, all other deposition parameters were constant, where scan speed and powder flow rate were selected as 1000 mm/min and 2.1 g/min, respectively. For the sample obtained at 275 W, obvious gas porosity with circular shapes was observed with a large variation of pore sizes. A significant proportion was in the medium to large range. The higher laser power at 485 W significantly reduced the frequency and average size of pores within the printed samples. Nearly all medium sized pores were eliminated, with only a few large pores and mostly small pores present. The reduction in porosity was confirmed by a pore area analysis. The porosity of the samples produced at 275 W was close to 10%, whereas for the samples produced at 485 W, the porosity was significantly reduced to just 1.7%, which was the lowest porosity of all samples analyzed in this study. A weighted histogram in Fig. 10 better illustrates the changes in pore size distribution after the laser power was increased from 275 W to 485 W. The pore size was calculated through equivalent circular area diameter [32]. The histogram was normalized by the whole cross-sectional area analyzed, so the height of each bar represented the area fraction of pores within the corresponding size range. It was seen from Fig. 10 that the fraction of pores of all size ranges reduced at the higher laser power of 485 W. Meanwhile, the reduction in the fraction of pores larger than 30 μm was more substantial than that for pores within other size ranges. It is worth noting that the elimination of large pore sizes will be beneficial for obtaining transparency. If pores could be reduced to a small enough size, high transparency would be possible even with a substantial number of pores present [23]. The reduction in large pore size at higher laser power can be attributed to a higher volumetric energy density of 116 J/mm3 at 485 W compared to 82.5 J/mm3 at 275 W. The increase in volumetric energy density was the primary reason for increased melt pool temperatures, longer solidification times, and reduced viscosity. These factors combined to allow more trapped gasses to combine and escape to the melt surface before solidification, thus mostly eliminating larger pores. Due 6

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Fig. 8. Top layers of thin wall structure with (A) No additional laser heat treatment (B) Laser heat treatment by ten additional laser scanning passes at a scan speed of 1000 mm/min with no powder feeding.

to increased melt size, remelting of deposited samples also helped remove the trapped pores in previously deposited layers. It should be noted that samples produced at high laser power were subject to decreased melt stability and resulted in irregular shaped structures. Fig. 9(C) shows a specimen fabricated at 275 W. The shape of this specimen was uniform along the build direction due to a relatively good print resolution. On the other hand, the sample printed at 485 W in Fig. 9(D) exhibited an irregular shape. Increased melt temperatures at the higher laser powers resulted in reduced melt viscosity, which lead to substantial melt pool spreading and sagging. This resulted in a significantly wider, misshapen structure in Fig. 9(D). The loss of print resolution seen at high laser power posed additional challenges in printing near net shape complex optical components.

3.2.3. Effects of pulsed laser mode on residual porosity Besides a variation of laser power in the CW mode discussed above, the effects of laser operation modes, i.e., CW mode and pulsed mode, were also studied here with respect to porosity and densification. Electromagnetic radiation during laser direct deposition exerts a small amount of pressure when interacting with a material. When lasing a substance in pulsed mode, there is a constant fluctuation of incident radiation, and thus exerted pressure, which results in vibration with the same frequency as the pulsed laser beam. This vibration in the melt pool could aid in gas removal. Thus, laser direct deposition of spinel using pulsed laser mode was tested to determine what effect, if any, it would have on the resultant porosity of produced parts. The pulsed laser parameters, varying in pulse period and pulse frequency, are summarized in Table 3. In this study, average laser power was held constant at 275 W to have a direct comparison with the samples printed in the CW mode. Scan speed and powder flow rate were fixed at 1000 mm/min and 2.1 g/min, respectively. The measured average porosity and density were summarized in Fig. 11(A) and Fig. 11(B), respectively. Compared to the samples printed in the CW mode, no obvious reduction in porosity or increase in density was observed when the deposition was performed in the pulsed mode. While deposited in the pulsed mode, no clear trend in the variation of porosity or density was found with varied laser pulse width and frequency. The average porosity in Fig. 11(A) ranged between 9.5% and 12.4% with the largest value observed at a pulse width of 1.5 ms and a pulse frequency of 135 Hz. The smallest average density was also found for the samples printed at this intermediate pulsed mode in

Fig. 9. The effects of laser power: (A) Micrograph of sample cross-section fabricated at 275 W showing a high frequency of large pores, (B) Micrograph of sample cross-section produced at 485 W with greatly reduced pore size and frequency, (C) Thin wall structure printed at 275 W showing good printability and uniform layers, and (D) Thin wall structure exhibiting irregular shape and poor resolution due to a higher laser heat input at 485 W.

Table 3 Laser parameters for pulsed laser deposition of spinel ceramics.

Fig. 10. A weighted histogram showing the distribution of pore size for the samples printed at 275 W and 485 W.

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Laser Mode

Pulse Width (ms)

Pulse Frequency (Hz)

Average Power (W)

Pulsed Pulsed Pulsed Continuous

0.24 1.50 9.99 N/A

333 135 71 N/A

275 275 275 275

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Fig. 11. Results in comparison of laser CW and pulsed modes on (A) Porosity area, and (B) Sample density.

reduce the distance that gasses needed to travel before escaping from the melt pool, allowing more gasses removed before solidification and thus leading to the lower porosity observed at 0.9 g/min compared to that at 2.1 g/min. The pore size distribution was further analyzed for the samples with the maximum and minimum porosity, which were printed at the powder flow rates of 2.1 g/min and 0.9 g/min, respectively. The obvious reduction in the average porosity seen in Fig. 12 was found to be attributed to the significantly reduction of the pores larger than 30 μm in Fig. 13. In contrast, it is interesting to note that at the lower powder flow rate, an obvious increase of smaller pores (< 30 μm) was even noticed. It was possible that the increased effective laser heat input at the powder flow rate of 0.9 g/min led to the vaporization of Al2O3–MgO melt pool, which often started as small pores. Further studies will be performed to analyze the melt pool temperature distribution and examine the composition change due to the potential vaporization process. The change in pore size distribution was even more obvious after decreasing the powder spot size from 5 mm to 1 mm. The reduction in spot size was achieved by replacing the powder delivery tube of 1.6 mm inner diameter with a 0.8 mm inner diameter tube. The corresponding powder flow rates were reduced from 2.1 g/min to 0.8 g/min. Despite an obvious reduction in average porosity in Fig. 14(A) and a significant reduction in pores larger than 30 μm in Fig. 14(B), additional, smaller pores were found for the samples printed at a powder spot size of 1 mm. This was also believed to be related to a similar vaporization process due to increased effective energy input at the reduced powder flow rate. Also, the decrease of powder spot size may further promote the vaporization. If the powder spot (5 mm) was larger than the laser spot (2.5 mm throughout this study), a significant amount of powder would

Fig. 11(B). The densities of samples, ranging between 92.5% and 94.1%, were within the uncertainty margin using different pulsed parameters, thus showing no apparent change in densification using pulsed mode over CW mode at the selected laser parameters. However, a recent study [33] suggested that the implementation of an ultrafast pulsed laser, i.e., femtosecond and picosecond pulses, may exhibit dramatically different localized melting, compared to the millisecond pulses used in this study, thus worth further investigation. 3.2.4. Effects of powder flow rate and powder spot size on residual porosity Powder flow rate was found to have a significant effect on residual porosity as seen in Fig. 12. Increasing powder flow rate consistently increased porosity and decreased the density of fabricated parts. Fig. 12(A) shows that the part porosity with a powder flow rate of 2.1 g/min had an average porosity of 9.8% while the samples produced with a powder flow rate of 0.9 g/min only had an average porosity of 5.2%. The measured density in Fig. 12(B) correlated well with the findings in the porosity analysis, where an increase in relative density was observed with a decrease in powder flow rate. Using a powder flow rate of 2.1 g/min the average relative density was 92.5%. Printed at a lower powder flow rate of 0.58 g/min, the average relative density increased to 97.6%. While the laser power and scan speeds were maintained constant in this study (at 275 W and 2000 mm/min, respectively), the effective energy input increased at a decreased powder flow rate [34]. Higher energy input for the volume of material deposited increased localized temperatures that reduced the viscosity of the melted spinel. This allowed for more of gas within the parts to combine and if the buoyant forces were sufficient, escape through the melt surface. Also, by reducing powder flow rate, layer thickness was reduced from 0.8 mm to 0.2 mm. The reduced layer thickness would

Fig. 12. Results from powder flow rate study on (A) Porosity area and (B) Sample density. 8

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that resulted during powder mixing. When large voids were present in the initial powder agglomerates, gasses were transported directly into the melt pool, which, depending on processing conditions, may not have the sufficient time to escape before the melted mixture solidified, thus increasing the porosity level [16]. Meanwhile, the reduction in melting point for the nano-MgO powders would also potentially contribute to the reduced porosity level. At a lower melting point, the nanoparticles would melt faster, better wetting the Al2O3–MgO powder mixtures while allowing more time for gasses to escape the melt before solidification. Fig. 16 shows a weighted histogram illustrating the fraction of different pore sizes within the samples printed using both powder compositions. It is clearly seen that the samples printed with nano-MgO particles had a significant change in pore size distribution. The percentage of large pores (> 30 μm) was significantly reduced from 8.3% with micron-MgO powder down to 1.8% with nano-sized MgO powder. In contrast, no obvious difference in the fraction of smaller pore sizes was observed between these two powder compositions. The fraction of pore size smaller than 10 μm even slightly increased for the samples printed with nano-sized MgO powder. This could be attributed to an increase in small pores in the agglomerates of Al2O3–MgO powder mixtures, where the nano-MgO powder was unable to infiltrate during the acetone removal and powder drying process. Another potential factor affecting the porosity was a difference in the compositions of the printed samples as two different initial MgO particle sizes were used in this study. Thus, the XRD patterns for the samples printed with both types of MgO powders, i.e., micron-MgO and nano-MgO powders in this study, were also measured as shown in Fig. 17. The results showed that only a single spinel phase was present for both types of samples. The detected peaks matched the reference peaks for stoichiometric magnesium aluminate spinel (MgAl2O4). This indicated that the spinel ceramics could be directly synthesized from the Al2O3–MgO powder mixtures via melt-grown methods of one-step laser direct deposition in this study. It should be noted that as the XRD analysis was directly performed on the polished bulk samples, the sample texturing might contribute to the difference in the detected peak intensity in Fig. 17. The results in Fig. 17 also showed a slight peak shift. This indicated that an alumina-rich single-phase spinel was possibly produced [35], and therefore a degree of vaporization may have occurred, leading to the formation of small pores observed in this study.

Fig. 13. A weighted histogram showing the change in the pore size distribution for the samples after decreasing the powder flow rate from 2.1 g/min to 0.9 g/ min.

flow outside the laser focus area with a lower energy input. On the other hand, if the powder spot (1 mm) became smaller than the laser spot, all the powders would be directly melted or even vaporized by the focused laser beam at a much higher energy input, thus increasing the chance of the formation of smaller pores due to vaporization.

3.2.5. Effects of initial powder particle size on residual porosity Using powder mixtures prepared with nano-particle MgO, a substantial porosity reduction (nearly 60%) was observed in Fig. 15(A), which shows a direct porosity comparison between samples produced using the two different MgO powders in this study. The laser power, scan speed, and powder flow rate were maintained at constant values of 275 W, 1000 mm/min, and 2.1 g/min, respectively. Fig. 15(B) shows a representative cross-section of the samples fabricated with average MgO particle size < 44 μm. This sample had relatively large pores throughout the cross-section. In comparison, the cross-section of the samples produced with average MgO particle size < 50 nm in Fig. 15(C) showed an obvious reduction in pore size and frequency. The porosity reduction could be attributed to the fact that the nanoparticles better filled in large voids in the Al2O3–MgO powder mixtures. In the micron-MgO powder mixture, additional gasses were trapped within powder agglomerates (due to poorer packing efficiency)

4. Conclusions In this work, the effects of processing conditions on residual porosity was investigated in additively manufactured spinel ceramics by laser direct deposition. Both processing parameters and powder

Fig. 14. Results from powder spot size study on (A) Average porosity, and (B) Pore size distribution. 9

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Fig. 15. Results using different initial MgO particle sizes: (A) Measured porosity levels, (B) Cross-sectional micrograph of thin wall produced with MgO average particle size < 44 μm showing large pores, and (C) Cross-sectional micrograph fabricated with MgO average particle size < 50 nm exhibiting significantly reduced pore size and frequency.

2000 mm/min followed by a subsequent decrease at higher scan speeds. Increased laser power was shown to dramatically reduce porosity in bulk ceramics through increased melt temperatures and longer melt durations. A similar porosity reduction was also observed at reduced powder flow rate and powder spot size due to increased localized temperatures and decreased layer thickness, both of which aided in gas removal before solidification. It was demonstrated that using nanoparticle MgO powder as a precursor, a significant reduction in residual porosity was achieved due to the reduced large pores present in powder agglomerates for fabrication. With no obvious change of smaller pore proportion, the average porosity reduction was mainly due to the significant reduced fraction of large pores (> 30 μm). Laser CW mode and millisecond pulsed mode showed no obvious effects on porosity and densification. Through the detailed studies of porosity and densification of laser direct deposited spinel, a better understanding of the pore removal process can further help control pore sizes and fractions during AM of spinel to obtain high-quality optical ceramic components. Declaration of competing interest None.

Fig. 16. Histogram showing pore size distribution within the samples printed with MgO initial average particle size < 44 μm and < 50 nm.

Acknowledgements The authors would like to extend a special thanks to Dr. Jeremy Watts for advice and continued assistance with sample preparation for microstructural analysis. We would also like to thank Joe Atria at Almatis for generously supplying the alumina powder used in this study. The assistance of Aditya Thakur in taking scanning electron micrographs of the powder mixtures is also much appreciated. References [1] T. Benitez, S.Y. Gómez, A.P.N. de Oliveira, N. Travitzky, D. Hotza, Transparent ceramic and glass-ceramic materials for armor applications, Ceram. Int. 43 (16) (2017) 13031–13046. [2] A.A. DiGiovanni, L. Fehrenbacher, D.W. Roy, Hard transparent domes and windows from magnesium aluminate spinel, Wind. Dome Technol. Mater. IX 5786 (2005) 56. [3] F. Klocke, O. Dambon, T. Bletek, T. Höche, F. Naumann, T. Hutzler, Surface integrity in ultra-precision grinding of transparent ceramics, Procedia CIRP 71 (2018) 177–180. [4] L. Lallemant, G. Fantozzi, V. Garnier, G. Bonnefont, Transparent polycrystalline alumina obtained by SPS: green bodies processing effect, J. Eur. Ceram. Soc. 32 (11) (2012) 2909–2915. [5] D.C. Harris, et al., Refractive index of infrared- transparent polycrystalline alumina, 56 (7) (2019). [6] F. Yan, W. Xiong, E.J. Faierson, Grain structure control of additively manufactured metallic materials, Materials 10 (11) (2017) 1260. [7] J. Wilkes, Y.C. Hagedorn, W. Meiners, K. Wissenbach, Additive manufacturing of ZrO2-Al2O3ceramic components by selective laser melting, Rapid Prototyp. J. 19 (1) (2013) 51–57. [8] M. Trunec, K. Maca, R. Chmelik, Polycrystalline alumina ceramics doped with

Fig. 17. XRD pattern for the samples printed with two types of MgO powders: micron-MgO with average particle sizes less than 44 μm and nano-MgO with average particle sizes less than 50 nm.

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