- Email: [email protected]
On the microstructure and properties of an advanced cemented carbide system processed by selective laser melting
Accepted Manuscript On the microstructure and properties of an advanced cemented carbide system processed by selective laser melting Chen-Wei Li, Kai-Chun Chang, An-Chou Yeh PII:
S0925-8388(18)34743-1
DOI:
https://doi.org/10.1016/j.jallcom.2018.12.187
Reference:
JALCOM 48828
To appear in:
Journal of Alloys and Compounds
Received Date: 11 October 2018 Revised Date:
11 December 2018
Accepted Date: 13 December 2018
Please cite this article as: C.-W. Li, K.-C. Chang, A.-C. Yeh, On the microstructure and properties of an advanced cemented carbide system processed by selective laser melting, Journal of Alloys and Compounds (2019), doi: https://doi.org/10.1016/j.jallcom.2018.12.187. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
ACCEPTED MANUSCRIPT
On the microstructure and properties of an advanced cemented carbide system processed by selective laser melting Chen-Wei Lia, Kai-Chun Changa, An-Chou Yeha, b* Department of Materials Science and Engineering, National Tsing Hua University,
RI PT
a
Hsinchu City 30013, Taiwan, ROC b
High Entropy Materials Center, National Tsing Hua University, Hsinchu City 30013,
Taiwan, ROC
SC
*Corresponding author: [email protected]
M AN U
Abstract
This article presents a study on the influence of selective laser melting (SLM) process on microstructure and property of a cemented carbide system containing high entropy alloy. Analysis along the building direction indicated variation of chemical composition and microstructure, and this was influenced by two effects, firstly the dilution effect due to elemental diffusion from the baseplate and secondly the
TE D
elements evaporation caused by high-power laser. At the lower half of the specimen, high fraction of η-carbide formed near the level of baseplate, and there were chemical gradients of major binder elements along the building direction. At the upper half of the specimen, there were relatively less variation in chemical composition and more
EP
homogeneously distributed phases including WC, W2C, η-carbide and FCC metal binder. The hardness of the lower half specimen varied from 711.7 HV1 (bottom of
AC C
the specimen) to 1178.6 HV1 at 1 mm height. For the upper half of the specimen, hardness values could range from 1306.8 HV1 to 1413.4 HV1 and fracture toughness varied from 9.74 MPa m1/2 to 13.29 MPa m1/2. Keywords: additive manufacturing; selective laser melting; cemented carbide; microstructural characterization
1
ACCEPTED MANUSCRIPT
1. Introduction Cemented carbide consisting of hard tungsten carbide particles and ductile metal binder phase can possess extraordinary hardness, wear-resistance, and toughness. Tools made of cemented carbide have been widely used in industrial applications, such as machining, cutting or drilling [1-3]; however, traditional fabrication route
RI PT
through liquid phase sintering can only make mold with symmetrical and simple geometry [4, 5]. Thus, in order to manufacture cemented carbide tools with complex geometry to meet advanced industrial requirements, scientists and engineers are exploring possibilities to employ additive manufacturing (AM) processes [5-12]. For the fabrication of cemented carbides, there are various AM processes that have been
SC
attempted and they can be classified by the need for post-sintering process. For those techniques requiring post-sintering, including 3D gel-printing [5] and Thermoplastic
M AN U
3D Printing [5, 6], the cemented carbide raw materials are blended with organic binders as a slurry for shape-forming. The slurry is then dispatched by a nozzle to form green parts that need to be sintered to remove organic ingredients as well as to obtain the designed geometries. On the other hand, laser engineered net shaping [7, 8], electron beam melting [9], selective laser sintering/melting [13] are those processes without the need for post-sintering process; thus, the powders are free from organic
TE D
ingredients and can be directly consolidated by laser or electron beam via sinter/melting to form near-net-shape parts. Among all techniques, selective laser melting (SLM) and selective laser sintering (SLS) processes have received the most attention. Previous studies have utilized powders with high fraction of carbide
EP
particles (> 80 wt%) [14-16]; bronze-infiltrated WC-9Co composite had been fabricated by SLS process, although the wear resistance of the as-built mold could be
AC C
comparable to its counterpart made by tool steel, the infiltrated WC-Co composite could not be 100% dense due to the nature of SLS process. To further increase the density, partial melting of WC by laser with higher power had been suggested [14]. Uhlmann E. et al. [15] utilized agglomerated and pre-sintered powders for SLM process under various laser parameters; they unveiled that the relative density of cemented carbide parts could only be improved at the cost of much loss of Co by evaporation. They also concluded a conflict between the optimization of the relative density of SLM cemented carbide parts and the remaining content of Co, which could influence the toughness of the built parts. Domashenkov A. et al. [16] adopted conventional and nano-structured WC-12Co composite powders for SLM process and 2
ACCEPTED MANUSCRIPT investigated the influence of different powders on the microstructure and mechanical properties of the as-built cemented carbide; they failed to achieve 100 % dense part; however, they found that cemented carbide parts made from nano-phase powder could possess higher homogeneity in microstructure. Among the aforementioned studies, Co has been the most popular choice as the binder material due to its outstanding intrinsic
RI PT
properties. Furthermore, its moderate solubility in WC not only results in an excellent wettability between Co and WC, but also makes it easier to retain the carbon content during sintering [17, 18]; however, due to the high cost and high toxicity of Co [19], substitution of Co as the binder metal has been a subject of interest for decades [20-24]. Our previous work used Ni as the binder metal for WC-W2C, and a dense
SC
microstructure could be achieved by SLM process [25]. Recently, multi-element alloy such as high entropy alloy (HEA) has been used as the binder metal to fabricate
M AN U
cemented carbide by liquid phase sintering process, and inhabitation of the growth of WC particles during the sintering process by the complex alloy system has been reported [26, 27]. For laser-based processes such as laser welding, laser cladding and SLM/SLS, the properties of the built can be affected by the inter-diffusion or dilution between the baseplate material and the deposited material [28]. However, the proposed sluggish diffusion of the HEA might hinder interdiffusion in cemented
TE D
carbide systems [29, 30] during the SLM process. Previous work tried to laser-clad several layers of AlCoCrCuFeNi HEA onto a Mg-substrate, with proper process parameters, no serious dilution occurred on the HEA composition, which preserved the designed properties of the binder alloy [31]; however, to the best of the authors’
EP
knowledge, to date there is no reported study on attempting to build cemented carbide with high entropy alloy binder by SLM process. This study aims to examine the use
AC C
of HEA as a binder material for cemented carbide by investigating the effect of SLM process on chemical composition of the built and properties; this study can provide a reference for future development of cemented carbide processed by SLM.
2. Experimental procedure A powder mixture consisting of 80 wt.% of spherical cast tungsten carbide powder (TEKMAT WC-45, TEKNA Advanced Materials Inc., Canada) and 20 wt.% of gas-atomized multi-element NiAlCoCrCuFe high entropy alloy powder was utilized in this study. The microstructure of both powders were given in Fig. 1. The binder alloy NiAlCoCrCuFe utilized Ni, Fe, Co as the major elements (> 19 wt.%) 3
ACCEPTED MANUSCRIPT and Al, Cr, Cu as the minor elements (< 12 wt.%). The addition of Ni and Fe aimed to decrease the use of Co; Cr, Al and Cu were for solid solutions and oxidation resistance [27, 32]. The nominal compositions of both forms of powders are given in Table 1, and the phase constitutions of both powders were checked by XRD as given in Fig. 2. The NiAlCoCrCuFe alloy powder comprised a single FCC phase.
RI PT
cast tungsten carbide powder consisted of two phases, i.e. WC and W2C, and the The mean particle sizes of carbide powder and NiAlCoCrCuFe alloy powder were measured by Laser Diffraction Particle Size Analyzer (Coulter LS230, Beckman Coulter Inc., U.S.) as 27.69 µm and 47.88 µm, respectively. The Carr’s
SC
compressibility index of the powder mixture was 3.3 %, which indicated an excellent flowability for powders and its suitability for powder-bed type SLM process.
M AN U
The SLM machine utilized in this experiment was a custom-made setup equipped with an ytterbium fiber laser (λ = 1070, YLR-500-SM-AC, IPG Photonics Co., U.S.). The diameter of the laser spot was 57 µm on the fixed focal plane. The oxygen level of the working chamber was kept below 1500 ppm with a protective argon circulation system to prevent oxidation. A baseplate made of Invar alloy (36 wt.% Ni, 64 wt.% Fe) was used due to its low value of thermal expansion. Cemented
TE D
carbide specimens were built on top of the baseplate as square blocks (15 mm × 15 mm × 3 mm) by a set of optimized laser parameters (140 W of laser power, 90 mm/s of scan speed, 115 µm of hatch distance, and 40 µm of layer thickness). As-built cemented carbide specimens were sectioned from the baseplate and cut
EP
transversely by wire-cut electrical discharge machining to reveal the cross-sectional microstructures. Samples were mounted, mechanically ground and polished by
AC C
diamond grinding disk and diamond polishing suspensions down to 3 µm, followed by a final 0.05 µm alumina suspension polishing. The as-polished microstructure was then etched by Murakami’s reagent (1 g KOH, 1 g K3[Fe(CN)6], 100 ml H2O) for 12 seconds, therefore the present phases could be identified based on their different reaction rates to Murakami’s reagent [33]. For example, W2C should be rapidly etched and shown as dark-color; on the other hand, WC and η-phase were more resistant to Murakami’s reagent and remained to be lighter in color contrast. Microstructure observations were conducted by BX-51 optical microscope (OM, Olympus Co. Ltd., Japan), SU-8010 field-emission scanning electron microscope (FE-SEM, Hitachi Ltd., Tokyo, Japan) and JSM-IT100 SEM (JEOL Ltd., Tokyo, 4
ACCEPTED MANUSCRIPT Japan). Analyses on chemical composition were performed by JSM-IT100 SEM with its attached INCA x-act energy dispersive x-ray spectrometer (EDS, Oxford Instruments plc, UK). Phases were identified by D2 Phaser X-ray diffractometer (XRD, Bruker Co. Ltd., U.S.) with Cu Kα radiation (λ = 1.54 Å), and the measurement was set to scan
RI PT
with a fixed step size of 0.02º. Samples were ground and polished for accurate phase identification. For EBSD analysis, JEOL SM-09010 Cross-Section Polisher (JEOL, Ltd., Japan) was utilized to prepare ultra-fine sample surfaces. The surface polishing was carried out by argon ion under 6 kV of accelerating voltage for 16 h. Crystallographic information for EBSD characterization was obtained from PDF-4+
SC
2015 database (International Center of Diffraction Data, U.S.).
Vickers hardness test was carried out by model HM-115 micro vickers hardness
M AN U
testing machine (Mitutoyo Co., Japan). The Vickers hardness value was calculated according to HV = 1.8544 × f/d2, where f (kgf) is the load force and d (mm) is the mean diagonal length of the indentation. For the hardness profile, a load force of 1 kgf and a load duration of 12 seconds were set. From the original surface level of the baseplate (z = 0), hardness test was conducted every 0.25 mm along the building direction and 5 impressions were made for an average value. For the indentation
TE D
toughness test, a load force of 10 kgf was applied to form adequate indentation impressions and Palmqvist cracks at the corners. The toughness test was carried out at height level of z = 1 mm, 1.5 mm, 2.0 mm and 2.5 mm. Four indentation impressions were made for an average value. The fracture toughness (KIC) was determined based
EP
on formula (1), which was suggested by Niihara [34], Warren and Matzke [35]: (1)
AC C
KIC = β(HW)1/2 MPa m1/2,
where β is a constant and generally equals to 8.89 × 10-2 for tungsten carbide composite [36]; H is the hardness expressed in MPa; W is the crack resistance expressed as (P/Σl), which is the load force (P) over the total length (Σl) of all Palmqvist cracks originated from the corners of the Vickers hardness impression. Phase diagrams were simulated by CALPHAD-based Thermo-Calc software with TCFE7 database [37]. Phase diagrams were constructed according to the chemical composition measured by SEM-EDS. The temperature range (y-axis) and the carbon fraction (x-axis) were set from 1273 K to 3273 K and 1 wt.% to 3.5 wt.%, 5
ACCEPTED MANUSCRIPT
RI PT
respectively to cover the range of interest.
Fig. 1. Microstructure of (a) cast WC/W2C tungsten carbide powder and (b)
SC
gas-atomized NiAlCoCrCuFe alloy powder
M AN U
Table 1. Nominal compositions of WC-45 cast tungsten carbide powder and NiAlCoCrCuFe atomized alloy powder Binder
Ni
Al
Co
Cr
Cu
Fe
wt.%
39.36
4.89
20.29
4.71
11.52
19.23
Element
W
C
Fe
Co
Ni
Others
wt.%
Balance
3.0-4.1
< 0.3
< 0.1
< 0.1
< 0.1
AC C
EP
TE D
WC-45
Element
6
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
Fig. 2. XRD profile of (a) the cast tungsten carbide powder shows a constitution of WC and W2C phase, and (b) the NiAlCoCrCuFe alloy powder consists of a single
TE D
FCC phase.
3. Results and Analysis
3.1. Profile of chemical composition
The as-built SLM cemented carbide and its cross-section micrograph are shown
EP
in Fig. 3 and Fig. 4. The density and the fraction of porosity within were determined by Archimedes’ method as 13.46 g/mm3 and 0.64 %, respectively. Dense SLM
AC C
cemented carbide specimens have been fabricated in this work. The SLM process could induce melting of the baseplate and causing dilution effect on the chemical composition of the cemented carbide. Therefore, the chemical compositions at different heights of the built were analyzed by EDS and listed in Table 2 as well as plotted in Fig. 5 (carbon was excluded due to inaccuracy); about 0.45 mm beneath the baseplate surface possessed the initial Invar alloy composition. At the lower part of the built (from z = 0 mm to z = 1 mm), the weight fraction of Fe and Ni surpassed those of the nominal quantities and compositional gradients of Fe, Ni and W could be identified; however, in the upper part of the specimen (from z = 1 mm to the top surface), the fractions of major binder elements (Fe, Co, Ni) became less than the 7
ACCEPTED MANUSCRIPT nominal composition and remained constant. The measured/nominal ratios of major binder elements decrease from 2.35 (z = 0 mm) to 0.91 (z = 1 mm). Therefore, the lower part of the sample (from z = 0 mm to z = 1 mm) and the upper part of the sample (from z = 1 mm to the top surface) were termed as the region (measured/nominal ratio ≥ 0.91) and the region
(measured/nominal ratio ≤ 0.91),
M AN U
SC
RI PT
respectively, in this study.
AC C
EP
TE D
Fig. 3. As-built SLM cemented carbide specimen
8
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
Fig. 4. Optical micrograph of the cross-section of the as-built cemented carbide and the Vickers hardness impressions (load force = 1 kgf) along the building direction.
9
ACCEPTED MANUSCRIPT Table 2. Chemical compositions along the building direction by EDS analysis (wt.%) Major Alloy Element Site
Height (mm)
Minor Alloy Element
W
Fe
Co
Ni
Cr
Cu
Al
Measured/ Nominal
-
79.41
3.96
4.18
8.10
0.97
2.37
1.01
-
1
-0.45
2.43 ±
56.70 ±
0±
39.29 ±
0±
1.59 ±
0±
-
0.46
1.09
0.53
0.96
0.19
1.24
1.70
6
1.65 2.80
1.00
0.30
0.85
0.29
69.06 ± 10.58 ± 5.01 ± 12.64 ± 0.78 ± 1.19 ± 0.74 ± 1.37
0.66
0.56
0.81
81.18 ±
4.26 ±
3.87 ±
7.65 ±
1.25
0.50
0.57
0.72
85.6 ±
3.64 ±
3.44 ±
5.59 ±
1.13
0.46
0.52
0.64
86.33 ±
3.12 ±
3.44 ±
5.45 ±
1.21
0.50
0.55
0.68
0.31
0.64
0.26
0.67 ± 1.62 ± 0.76 ± 0.35
0.67 0.60 0.65
0.70
0.25
0.57 ± 0.87 ± 0.21 ± 0.36
0.91
0.27
0.62 ± 0.67 ± 0.44 ± 0.33
1.50
0.66
0.27
TE D
5
1.00
0.56
EP
4
0.44
1.08
2.35
AC C
3
0.14
51.66 ± 25.93 ± 2.36 ± 18.04 ± 0.17 ± 0.62 ± 1.21 ±
SC
0
M AN U
2
RI PT
Nom.
Fig. 5. Chemical composition profile of major binder elements and tungsten along the building direction.
10
ACCEPTED MANUSCRIPT 3.2. Microstructure Characterization To analyze the microstructure within region
(z = 0 - 1 mm), cemented carbide
specimen connected to a thin piece of baseplate was prepared; the cross-section micrograph is shown in Fig. 6(a). According to the back-scattered electron image shown in Fig. 6(b), deeply-penetrated molten pools (up to 300 µm in depth) beneath
RI PT
the original baseplate surface could be observed. Some un-melted tungsten carbide powders (indicated by red arrows) sank below the baseplate surface due to higher density within the molten pools. The applied high-power laser has led to melting and mixing of elements from carbide powders, alloy powders and baseplate alloy. η-carbides (M2W4C, M3W3C), W2C and FCC binder phase were the primary phases
SC
as characterized by XRD (Fig. 7). Due to the dilution of Fe and Ni from the baseplate, η-carbides formation were promoted [20]. High fraction of η-carbides was also
M AN U
identified by EBSD near the baseplate surface level (z = 0 mm) as shown in Fig. 8(a) and (b). Based on the image analysis results, an average volume fraction of 55.03 %
EP
TE D
of η-carbide and 44.97 % of metal binder phase were present at z = 0 mm.
AC C
Fig. 6. (a) OM and (b) SEM back-scattered micrographs of the region baseplate surface.
11
near the
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
Fig. 7. Four phases including WC, W2C, η-carbide and FCC metal binder could be identified by XRD within the SLM cemented carbide specimen. Phase constitutions , the region
, and the molten pools (top of the specimen)
TE D
varied between the region
AC C
EP
along the building direction.
Fig. 8. (a) SEM image and (b) EBSD phase mapping result identified the presence of η-carbide within the region
near the baseplate surface. (Yellow: M2W4C, Green:
M3W3C, Red: FCC, Blue: WC, Black: non-indexed point) At region
, microstructure contained faceted WC precipitates embedded in the 12
ACCEPTED MANUSCRIPT matrix, Fig. 9(a), and it shows the matrix featuring a W2C/metal dendritic structure labeled as site 1 in Fig. 9(b). The back-scattered electron image in Fig. 10 shows higher-magnified microstructure, which is at the same region as shown in sites 2 and 3 within Fig. 9(b); except for the light-grey W2C dendrite and the black inter-dendritic metal binder, a dark-grey phase formed heterogeneously and substitute some fractions
RI PT
of metal binder phase. The chemical composition of different phases (sites labeled as A, B, C, D and E) were checked by SEM-EDS and listed in Table 3. Comparing to the XRD analysis results, Fig. 7, site A and C could be identified as W2C phase for its high fraction of tungsten; sites B and D could be characterized as η-carbide for its higher fraction of Fe and Ni elements; site E could be identified as the FCC metal contained 13.14 % of WC,
SC
binder phase. Based on the image analysis results, region
39.61 % of W2C, 19.05 % of FCC binder phase and 28.20 % of η-carbide (at z = 2.0
TE D
M AN U
mm).
Fig. 9. Optical micrographs of the region
microstructure (a) as-polished and (b)
AC C
EP
etched by Murakami’s reagent
Fig. 10. Back-scattered electron image showing SEM micrographs of the region W2C are A and C, η-carbides are B and D, E is the metal binder. 13
.
ACCEPTED MANUSCRIPT Table 3. Chemical composition of different phases as labeled in Fig. 10 (in wt.%) Major Alloy Elements
Minor Alloy Elements
W
Fe
Co
Ni
Cr
Cu
Al
A
96.58 ±
1.06 ±
0.43 ±
1.19 ±
0.66 ±
0.08 ±
0 ± 0.11
0.49
0.19
0.21
0.25
0.17
0.29
82.90 ±
4.45 ±
4.59 ±
6.23 ±
0.85 ±
0.60 ±
0.38 ±
0.57
0.24
0.28
0.32
0.16
0.29
0.12
97.10 ±
1.01 ±
0.53 ±
0.76 ±
0.59 ±
0 ± 0.28
0 ± 0.11
0.39
0.19
0.20
0.24
0.16
84.19 ±
3.61 ±
4.19 ±
6.21 ±
0.56
0.23
0.27
0.32
47.04 ±
9.33 ±
9.08 ±
25.42 ±
0.83
0.30
0.34
0.54
D E
SC
C
0.84 ±
0.68 ±
0.29 ±
0.16
0.29
0.12
0.60 ±
4.81 ±
3.73 ±
0.15
0.37
0.17
M AN U
B
RI PT
Site
The micrograph given in Fig. 11(a) shows the top surface region of the built, which was the last solidified volume of the specimen, featuring keyhole-shaped molten pools at the top of region
. Fig. 11(b) shows the W2C/metal dendritic
TE D
colonies within the molten pools. Un-melted carbide powders were hardly seen in this region, which indicates that adopted high-power laser could fully melt both tungsten carbide powders and alloy powders within the molten pools. The depth of a molten pool could be up to 400 µm, which was 10 times the thickness of a single powder
EP
layer (40 µm) during the spreading process. Fig. 12 shows microstructure at the boundary between the molten pools and the
AC C
inter-pool area. Fine W2C/metal dendritic structure could be identified in the molten pool (left-hand side), and the inter-pool area was heat-affected due to overlapping of laser scanning and showed coarser W2C dendrites and WC precipitates (right-hand side); however, η-carbide was absent in the molten pools and the inter-pool area. Based on the image analysis results, the volume fraction of the dendritic W2C phase and the FCC binder phase were 87.30 % and 12.70 % in average at z = 2.5 mm, respectively.
14
RI PT
ACCEPTED MANUSCRIPT
EP
TE D
M AN U
SC
Fig. 11. (a) as-polished and (b) etched optical micrographs of the molten pool area.
AC C
Fig. 12. Microstructure at the boundary of a molten pool and the inter-pool area. Coarsened WC precipitates and W2C dendrites can be identified in the inter-pool area (right-hand side), while the molten pool remains a fine dendritic structure (left-hand side).
Vickers hardness test was carried out on the cross-section of the specimen along
the building direction, and the test results were drawn as a hardness profile shown in Fig. 13. In the region
, a gradient value of hardness is present and the average
hardness increases from 711.7 HV1 at z = 0 mm to 1178.6 HV1 at z = 1 mm. As for the region
, hardness values vary from 1306.8 HV1 (z = 1.25 mm) to 1413.4 HV1 (z
= 2.50 mm). On the other hand, indentation toughness test was also conducted along 15
ACCEPTED MANUSCRIPT the building direction, and the toughness profile is given in Fig. 13. Microstructure of the indentation impression and the radial cracks formed at the corners are given in Fig. 14. For the region
, its fracture toughness was too high to be determined by
indentation fracture method since no crack formed during the test. The region possesses fracture toughness ranging from 13.29 MPa m1/2 (at z = 1.0 mm) to 9.74
TE D
M AN U
SC
RI PT
MPa m1/2 (at z = 2.5 mm).
Fig. 13. Profile of Vickers hardness and indentation fracture toughness at different
AC C
EP
height level along the building direction.
Fig. 14. Micrographs of Vickers' indentation impressions and induced surface Palmqvist cracks in region II (a) z = 1.0 mm, and (b) z = 2.5 mm (molten pool area). 16
ACCEPTED MANUSCRIPT
4. Discussion According to the composition profiles shown in Table 2 and Fig. 5, the variation of composition within region
and region
can be governed by two mechanisms,
i.e. evaporation of elements during laser scanning and dilution by elements diffused from baseplate Invar alloy, i.e. Ni and Fe. Although the molten pools beneath the top
RI PT
surface of the specimen were not affected by the dilution from the baseplate, deviation from the nominal compositions in this region indicates that high-power laser could cause composition changes due to evaporation of elements, and this phenomena has been reported previously [38]. Within region
(z = 0 – 1 mm) and along the building
direction, the increasing species were W, Co, Cr, Cu, and the decreasing species were
SC
Fe, Ni and Al. Diffusion of Fe and Ni toward the built resulted in an average volume fraction of 55.03 % of η-carbide and 44.97 % of metal binder phase at z = 0 mm. The
M AN U
fractions of Co, Cr and Cu were lower than their nominal fractions at z = 0, possibly due to inter-diffusion toward the baseplate. As the cemented carbide kept building up, the molten pools moved upwards and there was less dilution effect from the baseplate, thus the fractions of other elements, such as W, Co, Cr, Cu were higher than those of the lower regions. In region
(z = 1 – 2.8 mm), W was the only element with
increased contents, and all the other binder elements, e.g. Fe, Co, Ni, Cr, Cu, Al,
TE D
decreased along the building direction, this was caused by the evaporation elements [38, 39], and the increase in W was primarily resulted from the loss of the other elements through vaporization.
The magnitude of elemental contents variation within the region
EP
smaller comparing to those of region
was much
, Fig. 5. Therefore, region II exhibited
relatively more uniform microstructure containing around 13.14 % of WC, 39.61 % of
AC C
W2C, 19.05 % of FCC binder phase and 28.20 % of η-carbide. The measured/nominal weight fraction ratios of the metal binder elements derived from Table 2 can be utilized to quantify the loss for each element in Table 4. The loss in weight fraction for elements are different from each other. According to the work done by Khan and Debroy [40], the rates of element loss during laser welding depend on the vaporization rates of element within the molten pools. The vaporization rates of element can be determined by the following equation [41]: Ji = AcPi/(2πMiRT)1/2
17
(2)
ACCEPTED MANUSCRIPT where Ji (kg-mole/m2 s) is the vaporization rate, Ac is a dimensionless parameter, Pi (N/m2) is the partial pressure of element i, Mi (kg/kg-mole) is the molecular weight of element i, R (kg m2/ s2 kg-mole K) is the gas constant, and T (K) is the temperature. In this study, the vaporization rates of different elements can be proportional to Pi/(Mi)1/2. According to He X. et al.[42], the peak temperature of the molten pools
RI PT
often exceeds the boiling temperature of elements. And the dendritic W2C present in the molten pools indicates that temperature during SLM should reach higher than the melting point of W2C, i.e. 3003 K [43], therefore, the ratios of vaporization rates for each elements can be derived from literatures and listed in Table 4. The major binder elements, Fe, Co and Ni have similar boiling points and can be compared directly
SC
with each other. Among the three major binder elements, Ni possesses the highest evaporation rate and Co exhibits the least evaporation rate, these analyses agree well
M AN U
with composition measurements in Table 2. For the minor elements, such as Cr, Cu, Al, their vapor pressures cannot be accurately estimated at temperature close to 3200 K due to their relatively lower boiling points and insufficient data; however, it can be reasonably assumed that elemental evaporation loss should be higher for those minor elements having lower boiling points. Therefore, the Al content loss was the most point.
TE D
severe among all the binder elements, and this can be resulted from its lowest boiling On the other hand, for around z=0 in region I, high power laser induced key-hole mode melting and melted not only cemented carbide but also a part of baseplate Invar alloy. As the depth of molten pools (300 µm) was larger than the thickness of one
EP
single powder layer (40 µm), a portion of the baseplate as well as the deposited part could be repeatedly re-melted (for the first several layers) during the SLM process.
AC C
Rapid convection and inter-diffusion of elements driven by the concentration gradients took place within the molten pools induced by high power laser [44]. Thus, comparing to the region
, more pronounced compositional gradients of all binder
elements within region
can be attributed to the dilution effect. As for the carbon
content, due to the difficulty to measure accurate carbon content by SEM-EDS, the variation of carbon fraction within different regions is discussed in the next section with the aid of both microstructure examination and predicted phase diagrams.
18
ACCEPTED MANUSCRIPT
Table 4 Properties, vapor pressure and partial pressure of present elements W
Fe
Co
Ni
Cr
Cu
Al
183.84
55.845
58.933
58.693
51.996
63.546
26.982
Boiling Point [45]
5828 K
3134 K
3200 K
3186 K
Nominal Molar
0.537
0.088
0.088
0.17
1.09
0.79
0.82
0.67
< 1 Pa
1.346 ×
1.013 ×
1.207 ×
5
10 Pa
5
Atomic Weight
2944 K
2835 K
2792 K
0.023
0.046
0.046
0.59
0.37
0.21
1.013 ×
1.327 ×
1.045 ×
10 Pa
5
10 Pa
5
10 Pa
105 Pa
At 3200
[47]
At 2944
[48]
[49]
K (B.P. of
At 3200
K (B.P. of
At 2900
At 2800
Co)
K
Cr)
K
K
0.891 ×
2.052 ×
0.233 ×
0.610 ×
0.481 ×
104 Pa
104 Pa
104 Pa
104 Pa
104 Pa
104 Pa
0.591
0.433
1
0.114
0.297
0.234
d ratio of weight Pi (at 3200 K)
5
10 Pa [46] At 3200 K
Ji/JNi
1.184 ×
-
TE D
Pi = XiPi0
M AN U
fraction 0
SC
Fraction (Xi) Nominal/Measure
RI PT
[45]
There were distinct composition gradients within the region
, so compositions
EP
at three positions were chosen for phase diagram simulations, results are shown in Fig. 15(a), 15(b) and 15(c) corresponding to z = 0 mm, 0.44 mm, and 1.00 mm,
AC C
respectively. By contrast, there was no distinct composition gradients found in region , so the composition of the molten pool area beneath the top surface was used to predict phase diagram of region II, Fig. 15(d). At z = 0 mm, due to the absence of WC, the rapidly cooled molten materials
were likely entered a two-phases domain (liquid + η) as labeled in Fig. 15(a) to form η-carbide at around 1900 K. The remaining molten metal solidified afterwards to form the FCC matrix. At z = 0.44 mm, with less degree of dilution and a lack of WC phase, the solidification process might follow the red arrow shown in Fig. 15(b), where materials went through liquid + W + W2C before reaching the final FCC + η-carbide. At z = 1 mm, the solidifying molten materials entered the “liquid + W2C” domain and 19
ACCEPTED MANUSCRIPT then the “η + WC” domain, Fig. 15(c). For region
, because of the decreased
fraction of metal binder (mainly Fe and Ni), the matrix changed to a W2C/metal dendritic structure as shown in Fig. 9(b). According to the phase diagram in Fig. 15(d), W2C dendrite formed first and followed by rapid solidification of the remaining molten metals as inter-dendritic FCC metal phase (route 1). The W2C/metal dendritic
RI PT
structure was then repeatedly heated by the iteratively laser scanning, which led to precipitation and coarsening of faceted WC from the W2C dendrite. Furthermore, at temperature below 1900 K, η-carbide was more thermodynamically stable than W2C. Phase transformation then took place between the dendritic W2C and the
SC
inter-dendritic metal phase, indicated by the following the reaction (3):
(3)
M AN U
FCC + W2C → WC + η-carbide
There was no distinct microstructural gradient along the building direction within region
, it was observed that phase transformation took place mainly within
the heat-affecting zone around the molten pools. At locations away from the molten pools, phase transformation would be halted due to insufficient thermal energy. For the molten pool area near the top surface, phase transformation took place within the
TE D
“Liquid + W2C” and “Liquid + W2C + WC” domains labeled in Fig. 15(d). The molten pool area featured as-solidified microstructure, which follows the route 3 solidification path in Fig. 15(d); for inter-pool areas, they were slightly heat-affected by the neighboring molten pools and followed the route 2 path. WC phase
EP
precipitated and coarsened from the W2C dendrite and at the inter-pool area. η-carbide was hardly seen within the inter-pool area. According to Fig. 15(d), η-carbide is
AC C
thermodynamically-stable at temperature below 1900 K, therefore, outside temperature range and insufficient thermal energy can be two reasons for the absence of η-carbide within the inter-pool area. Based on the analysis above, the carbon content can be estimated from the phase
diagrams as 1 – 1.2 wt% at z = 0 mm, 1.25 – 1.5 wt% at z = 0.44 mm, 2.3 – 2.6 wt% at z = 1 mm for the region WC) for the region region
and 2.6 – 3.3 wt% (Liquid + W2C → liquid + W2C +
. The decrease of carbon fraction was more pronounced in the
, which could be resulted from the dilution effect as well as diffusion of
carbon into the baseplate by inter-diffusion; however, for the region
, due to the
limited dilution effect by the baseplate elements, the carbon content remained close to 20
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
the nominal content.
21
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
Fig. 15. Thermo-Calc phase diagram simulations for the (a) region (b) region
(at z = 0.44 mm) (c) region
(at z = 0 mm),
(at z = 1.00 mm) (d) region
mm).
22
(at z = 1.65
ACCEPTED MANUSCRIPT The mechanical properties of the SLM cemented carbide vary between different regions within the built and can be related to their microstructures. Comparing the chemical composition profile (Fig. 5) with the hardness profile (Fig. 13), within the region
, the decreasing trend of major binder elements (Fe, Ni) corresponds to the
increase in hardness due to increased fractions of carbides. For region
, the weight
which reflected the hardness value plateau in the region
RI PT
fraction of major binder elements remained at 13-14 wt.% along the building direction, . Since the hardness of
cemented carbide can be proportional to the volume fraction of hard carbide phase [50], with increasing fraction of carbides and decreasing fraction of binder phase could lead to increased hardness as expected. As for the molten pool area beneath the inferior to the average value of the region
SC
top surface, there were only a few WC precipitates, and their fracture toughness was . WC phase possesses better fracture
within region
M AN U
toughness than that of W2C phase [51], thus the evenly-distributed WC precipitates could improve the fracture toughness of this region.
Mechanical properties of commercial cemented carbides [52, 53] and those of the present SLM built are summarized in Table 5 for comparisons.; the binder fractions of the SLM built in Table 5 include 30 wt.% at the middle of the regionⅠ(z = 0.5 mm), 14 wt.% at the region Ⅱ (z = 2.0 mm) and 13 wt.% at the top molten pool
TE D
area (z = 2.5 mm). The properties of the region Ⅰ can be compared with that of the commercial cemented carbide containing 25 wt% Co, and the higher hardness of the SLM built could be attributed to the presence of dendritic W2C, which is harder than that of WC [54, 55]. On the other hand, the region Ⅱ and the top molten pool area can
EP
be comparable to that of commercial cemented carbide containing 12 wt% Co. Although the hardness was slightly higher for the SLM-built parts, the more brittle
AC C
W2C dendrite also yielded lesser fracture toughness [56], especially for the top molten pool area.
Based on the simulated phase diagrams as given in Fig. 15, the formation of
brittle W2C and η-carbide were inevitable in the present system. Full-melting of carbide powders resulted in the W2C dendrite; while carbon-deficiency of the system led to formation of η-carbide. Since W2C could decompose into WC and W at 1300 ℃ [43], post heat treatment may induce WC formation for better toughness. Furthermore, compensation of carbon content may also reduce W2C and η-carbide contents by forming WC [22]. Considering the fabrication sequence of the layer-by-layer SLM process, region 23
ACCEPTED MANUSCRIPT Ⅰ with high fraction of metal binder phase and high fracture toughness can serve as a transition part between the region Ⅱ and the baseplate. The region Ⅱ with constant chemical composition and stable mechanical properties should be used as the main body of the SLM cemented carbide. However, the last-solidified molten pool regions with inferior toughness performance should be removed to enhance the
RI PT
performance. Table 5 Comparison of properties of selected conventionally-sintered cemented carbides (Krupp-Widia, Essen) and the present SLM built [53] Conventionally-Sintered
(wt%) Binder (wt%)
25Co
88
75
87
86
70
12
25
13
14
30
1290
780
1330.2
1395.8
988.5
9.7
11.9
-
Vickers Hardness
SC
Carbides
Top Molten
12Co
Pool
Region Ⅱ
M AN U
Type
Present SLM Cemented Carbide
Region Ⅰ
(kg/mm ) Fracture Hardness 1/2
12.7
14.5
EP
(MPa m )
TE D
3
In summary, SLM process yielded a gradient composition and microstructure
AC C
due to dilution effect from the inter-diffusion between the built and the baseplate. Although the dilution effect could only influence certain height of the built, in this case 1mm, evaporation of elements could also alter the composition of the cemented carbide. Furthermore, laser scanning tracks resulted in molten-pools that produced uneven distribution of carbides. Mechanical properties of the cemented carbides built by SLM were affected by the SLM processing characteristics mentioned above, and the most affected part was region I influenced by dilution effect. Furthermore, the application of HEA composition as the binder in this work could not hinder the uneven growth of carbides during laser scanning overlaps, and it had been very difficult to control the composition of NiAlCoCrCuFe high entropy alloy during the 24
ACCEPTED MANUSCRIPT SLM process due to different element has different diffusion characteristic and evaporation rate. The future design of the binder alloy should reduce the difference in boiling temperatures of the alloying elements, so a better compositional control for the binder might be achieved during SLM process. According to Figure 15, regardless of the variation in binder elements content,
RI PT
whether it was diluted by nickel and iron due to interdiffusion or evaporation due to high power laser scanning, the solidification process all involved in the formation of either brittle η or W2C that could render low fracture toughness of the built. It is possible that a designed scanning strategy can control the solidification path in order to induce more WC formation; prolong post heat treatment might also induce more
SC
WC formation. Future studies will be needed in these areas and are currently in
M AN U
progress.
5. Conclusions
In the present study, 80 wt.% of spherical cast tungsten carbide powder mixed with 20 wt.% of gas-atomized NiAlCoCrCuFe high entropy alloy powder was used to fabricate solid block cemented carbide sample by selective laser melting process. The effect of SLM process on the microstructure and mechanical properties have been 1.
TE D
examined and discussed. The followings are conclusions: On the cross-section of the SLM cemented carbide specimen, chemical composition could vary along the building direction due to elemental dilution from the baseplate and evaporation of elements. It can be characterized into two (from z = 0 to z = 1 mm) and region
EP
regions, region
(from z = 1 mm to 3
mm (top surface)). In region
, which was enriched with Fe and Ni from the baseplate material and
AC C
2.
therefore deviated from the nominal composition significantly, up to 55.03 %
η-carbide were embedded in the metal binder matrix. Region
features
W2C/metal dendritic matrix, evenly-distributed WC precipitates and η-carbide
which evolved from W2C dendrite and inter-dendritic metal binder phase. As for
the last-solidified molten pool area at the top of the region
, up to 87.3 % of
dendritic W2C and no η-carbide could be found, since keyhole-shaped molten pools located just beneath the top surface of the specimen were the least-heat-affected area throughout the specimen. 3.
Region
possesses hardness values vary from 711.7 HV1 (z = 0 mm) to 1178.6 25
ACCEPTED MANUSCRIPT HV1 (z = 1 mm); region
possesses hardness values ranging from 1306.8 HV1
(z = 1.25 mm) to 1413.4 HV1 (z = 2.50 mm). Furthermore, the region possesses moderate toughness ranging from 9.74 MPa m1/2 to 13.29 MPa m1/2 and should be considered as the main body of SLM cemented carbide, however, the top molten pool area should also be removed as well as the region I to obtain
RI PT
part with more desirable mechanical properties.
Acknowledgements
Authors would like to thank Prof. Jien-Wei Yeh and Prof. Su-Jien Lin for the use of hardness testing machine, and Prof. Jeng-Gong Duh for the use of JEOL SM-09010
SC
Cross-Section Polisher. This work was supported by the Ministry of Science and Technology (MOST) [grant numbers: 107-2218-E-007-012, 106-3114-E-007-011 and
M AN U
107-2218-E-007-016]; and the “High Entropy Materials Center” from The Featured Areas Research Center Program within the framework of the Higher Education Sprout Project by the Ministry of Education (MOE) and MOST in Taiwan [grant number: MOST 107-3017-F-007-003].
TE D
Declarations of interest: None
References
[1] L.J. Prakash, Application of fine grained tungsten carbide based cemented
EP
carbides, Int. J. Refract. Met. Hard Mater., 13 (1995) 257-264. [2] H.-O. Andrén, Microstructure development during sintering and heat-treatment of
AC C
cemented carbides and cermets, Mater. Chem. Phys., 67 (2001) 209-213. [3] M.J. Kupczyk, J. Komolka, High durability of cutting insert edges made of nanocrystalline cemented carbides, Int. J. Refract. Met. Hard Mater., 49 (2015) 225-231.
[4] K.J.A. Brookes, 3D-printing style additive manufacturing for commercial hardmetals, Met. Powder Rep., 70 (2015) 137-140. [5] X. Zhang, Z. Guo, C. Chen, W. Yang, Additive manufacturing of WC-20Co components by 3D gel-printing, Int. J. Refract. Met. Hard Mater., 70 (2018) 215-223. [6] J.P. U. Scheithauer, S. Weingarten, E. Schwarzer, A. Vornberger, T. Moritz, A. Michaelis, Droplet-Based Additive Manufacturing of Hard Metal Components by 26
ACCEPTED MANUSCRIPT Thermoplastic 3D Printing (T3DP), J. Ceram. Sci. Technol., 8 (2016) 155-160. [7] V.K. Balla, S. Bose, A. Bandyopadhyay, Microstructure and wear properties of laser deposited WC–12%Co composites, Mater. Sci. Eng., A, 527 (2010) 6677-6682. [8] Y. Xiong, J.E. Smugeresky, L. Ajdelsztajn, J.M. Schoenung, Fabrication of WC– Co cermets by laser engineered net shaping, Mater. Sci. Eng., A, 493 (2008) 261-266.
RI PT
[9] H. Peng, C. Liu, H. Guo, Y. Yuan, S. Gong, H. Xu, Fabrication of WCp/NiBSi metal matrix composite by electron beam melting, Mater. Sci. Eng., A, 666 (2016) 320-323.
[10] C.X. Wang, T. Laoui, J. Bonse, P.J. Kruth, B. Lauwers, L. Froyen, Direct Int. J. Adv. Manuf. Technol., 19 (2002) 351-357.
SC
Selective Laser Sintering of Hard Metal Powders: Experimental Study and Simulation, [11] J.P. Kruth, X. Wang, T. Laoui, L. Froyen, Lasers and materials in selective laser
M AN U
sintering, Assembly Autom., 23 (2003) 357-371.
[12] K. Maeda, T.H.C. Childs, Laser sintering (SLS) of hard metal powders for abrasion resistant coatings, J. Mater. Process. Technol., 149 (2004) 609-615. [13] G. Zong, Y. Wu, N. Tran, I. Lee, D. Bourell, J. Beaman, H. Marcus, Direct selective laser sintering of high temperature materials, in:
1992 International Solid
Freeform Fabrication Symposium, 1992.
TE D
[14] S. Kumar, Manufacturing of WC–Co moulds using SLS machine, J. Mater. Process. Technol., 209 (2009) 3840-3848.
[15] E. Uhlmann, A. Bergmann, W. Gridin, Investigation on Additive Manufacturing of Tungsten Carbide-cobalt by Selective Laser Melting, Proc. CIRP, 35 (2015) 8-15.
EP
[16] A. Domashenkov, A. Borbély, I. Smurov, Structural modifications of WC/Co nanophased and conventional powders processed by selective laser melting, Mater.
AC C
Manuf. Process, 32 (2017) 93-100. [17] G.S. Upadhyaya, Cemented tungsten carbides: production, properties and testing, first ed., William Andrew, New York, 1998. [18] H.E. Exner, Physical and chemical nature of cemented carbides, Int. Mater. Rev., 24 (1979) 149-173.
[19] B. Gries, L.J. Prakash, Hard Materials 1: Acute Inhalation Toxicity by Contact Corrosion-The Case of WC-Co, in:
European Congress and Exhibition on Powder
Metallurgy. European PM Conference Proceedings, The European Powder Metallurgy Association, 2007, pp. 189-196. [20] B. Uhrenius, H. Pastor, E. Pauty, On the composition of Fe-Ni-Co-WC-based 27
ACCEPTED MANUSCRIPT cemented carbides, Int. J. Refract. Met. Hard Mater., 15 (1997) 139-149. [21] G. Gille, J. Bredthauer, B. Gries, B. Mende, W. Heinrich, Advanced and new grades of WC and binder powder – their properties and application, Int. J. Refract. Met. Hard Mater., 18 (2000) 87-102. [22] S. Norgren, J. García, A. Blomqvist, L. Yin, Trends in the P/M hard metal
RI PT
industry, Int. J. Refract. Met. Hard Mater., 48 (2015) 31-45. [23] W.D. Schubert, M. Fugger, B. Wittmann, R. Useldinger, Aspects of sintering of cemented carbides with Fe-based binders, Int. J. Refract. Met. Hard Mater., 49 (2015) 110-123.
[24] M. Tarraste, J. Kübarsepp, K. Juhani, A. Mere, M. Kolnes, M. Viljus, B. Maaten,
SC
Ferritic chromium steel as binder metal for WC cemented carbides, Int. J. Refract. Met. Hard Mater., 73 (2018) 183-191.
M AN U
[25] C.-W. Li, K.-C. Chang, A.-C. Yeh, J.-W. Yeh, S.-J. Lin, Microstructure characterization of cemented carbide fabricated by selective laser melting process, Int. J. Refract. Met. Hard Mater., 75 (2018) 225-233.
[26] C.-S. Chen, C.-C. Yang, H.-Y. Chai, J.-W. Yeh, J.L.H. Chau, Novel cermet material of WC/multi-element alloy, Int. J. Refract. Met. Hard Mater., 43 (2014) 200-204.
TE D
[27] B.L. Ezquerra, T.S. Biurrun, L.L. Cabezas, J.M. Sánchez Moreno, F.I. Lopez, R.M. Pampliega, Sintering of WC hardmetals with Ni-Co-Cr-Ti-Al multi-component alloys, Int. J. Refract. Met. Hard Mater., 77 (2018) 44-53. [28] S. Sun, Y. Durandet, M. Brandt, Parametric investigation of pulsed Nd: YAG
EP
laser cladding of stellite 6 on stainless steel, Surface and Coatings Technology, 194 (2005) 225-231.
AC C
[29] C.-W. Tsai, Y.-L. Chen, M.-H. Tsai, J.-W. Yeh, T.-T. Shun, S.-K. Chen, Deformation and annealing behaviors of high-entropy alloy Al0.5CoCrCuFeNi, J. Alloys Compd., 486 (2009) 427-435. [30] S.G. Ma, J.W. Qiao, Z.H. Wang, H.J. Yang, Y. Zhang, Microstructural features and tensile behaviors of the Al0.5CrCuFeNi2 high-entropy alloys by cold rolling and subsequent annealing, Materials & Design, 88 (2015) 1057-1062. [31] T.M. Yue, H. Xie, X. Lin, H.O. Yang, G.H. Meng, Solidification behaviour in laser cladding of AlCoCrCuFeNi high-entropy alloy on magnesium substrates, J. Alloys Compd., 587 (2014) 588-593. [32] C.-J. Tong, M.-R. Chen, J.-W. Yeh, S.-J. Lin, S.-K. Chen, T.-T. Shun, S.-Y. 28
ACCEPTED MANUSCRIPT Chang, Mechanical performance of the AlxCoCrCuFeNi high-entropy alloy system with multiprincipal elements, Metall. Mater. Trans. A, 36 (2005) 1263-1271. [33] J.R. Davis, ASM specialty handbook: tool materials, ASM international, 1995. [34] K. Niihara, A fracture mechanics analysis of indentation-induced Palmqvist crack in ceramics, J. Mater. Sci. Lett., 2 (1983) 221-223.
RI PT
[35] R. Viswanadham, Science of hard materials, first ed., Springer Science & Business Media, New York, 1983.
[36] D.K. Shetty, I.G. Wright, P.N. Mincer, A.H. Clauer, Indentation fracture of WC-Co cermets, J. Mater. Sci., 20 (1985) 1873-1882.
[37] Thermo-Calc Software TCFE7 Steels/Fe-alloys database.
SC
[38] W.E. King, H.D. Barth, V.M. Castillo, G.F. Gallegos, J.W. Gibbs, D.E. Hahn, C. Kamath, A.M. Rubenchik, Observation of keyhole-mode laser melting in laser
M AN U
powder-bed fusion additive manufacturing, J. Mater. Process. Technol., 214 (2014) 2915-2925.
[39] V.-P. Matilainen, H. Piili, A. Salminen, O. Nyrhilä, Preliminary Investigation of Keyhole Phenomena
during Single
Layer
Fabrication
in
Laser Additive
Manufacturing of Stainless Steel, Phys. Procedia, 78 (2015) 377-387. [40] P.A.A. Khan, T. Debroy, Alloying element vaporization and weld pool (1984) 641-644.
TE D
temperature during laser welding of AlSl 202 stainless steel, Metall. Trans. B, 15 [41] S. Dushman, J. Pierce, Scientific foundations of vacuum technique, Phys. Today, 2 (1949) 27.
EP
[42] X. He, T. DebRoy, P. Fuerschbach, Alloying element vaporization during laser spot welding of stainless steel, J. Phys. D: Appl. Phys., 36 (2003) 3079.
AC C
[43] H.O. Pierson, 6 - Carbides of Group VI: Chromium, Molybdenum, and Tungsten Carbides, in:
Handbook of Refractory Carbides and Nitrides, William Andrew
Publishing, Westwood, NJ, 1996, pp. 100-117. [44] M.M. Collur, A. Paul, T. Debroy, Mechanism of alloying element vaporization during laser welding, Metall. Trans. B, 18 (1987) 733-740. [45] D.R. Lide, CRC handbook of chemistry and physics: a ready-reference book of chemical and physical data, CRC press, 1995. [46] P.D. Desai, Thermodynamic Properties of Iron and Silicon, J. Phys. Chem. Ref. Data, 15 (1986) 967-983. [47] P.D. Desai, Thermodynamic properties of nickel, Int. J. Thermophys., 8 (1987) 29
ACCEPTED MANUSCRIPT 763-780. [48] R. Hultgren, P.D. Desai, Selected thermodynamic values and phase diagrams for copper and some of its binary alloys, (1971). [49] P.D. Desai, Thermodynamic properties of aluminum, Int. J. Thermophys., 8 (1987) 621-638. Mater. Sci. Eng., 33 (1978) 125-133.
RI PT
[50] H.C. Lee, J. Gurland, Hardness and deformation of cemented tungsten carbide, [51] H. Taimatsu, S. Sugiyama, Y. Kodaira, Synthesis of W2C by reactive hot pressing and its mechanical properties, Mater. Trans., JIM, 49 (2008) 1256-1261. [52] V. Sarin, Comprehensive hard materials, Newnes, 2014.
SC
[53] G.S. Upadhyaya, Materials science of cemented carbides — an overview, Materials & Design, 22 (2001) 483-489.
M AN U
[54] M. Khechba, F. Hanini, R. Halimi, Study of structural and mechanical properties of tungsten carbides coatings, Nature & Technology, (2011) 9.
[55] E. Zeiler, S. Schwarz, S. Rosiwal, R. Singer, Structural changes of tungsten heating filaments during CVD of diamond, Materials Science and Engineering: A, 335 (2002) 236-245.
[56] L. Jacobs, M. Hyland, M. De Bonte, Study of the influence of microstructural
TE D
properties on the sliding-wear behavior of HVOF and HVAF sprayed WC-cermet
AC C
EP
coatings, Journal of thermal spray technology, 8 (1999) 125-132.
30
ACCEPTED MANUSCRIPT
Highlights
M AN U
SC
RI PT
Dilution effect, evaporation, and laser scanning tracks could affect the built Hardness and toughness varied due to variation in composition and microstructure
TE D
EP
A complex cemented carbide system was subjected to selective laser melting process The cemented carbide system contains a high entropy alloy Composition and microstructure of the cemented carbide were heterogeneous
AC C
S0925-8388(18)34743-1
DOI:
https://doi.org/10.1016/j.jallcom.2018.12.187
Reference:
JALCOM 48828
To appear in:
Journal of Alloys and Compounds
Received Date: 11 October 2018 Revised Date:
11 December 2018
Accepted Date: 13 December 2018
Please cite this article as: C.-W. Li, K.-C. Chang, A.-C. Yeh, On the microstructure and properties of an advanced cemented carbide system processed by selective laser melting, Journal of Alloys and Compounds (2019), doi: https://doi.org/10.1016/j.jallcom.2018.12.187. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
ACCEPTED MANUSCRIPT
On the microstructure and properties of an advanced cemented carbide system processed by selective laser melting Chen-Wei Lia, Kai-Chun Changa, An-Chou Yeha, b* Department of Materials Science and Engineering, National Tsing Hua University,
RI PT
a
Hsinchu City 30013, Taiwan, ROC b
High Entropy Materials Center, National Tsing Hua University, Hsinchu City 30013,
Taiwan, ROC
SC
*Corresponding author: [email protected]
M AN U
Abstract
This article presents a study on the influence of selective laser melting (SLM) process on microstructure and property of a cemented carbide system containing high entropy alloy. Analysis along the building direction indicated variation of chemical composition and microstructure, and this was influenced by two effects, firstly the dilution effect due to elemental diffusion from the baseplate and secondly the
TE D
elements evaporation caused by high-power laser. At the lower half of the specimen, high fraction of η-carbide formed near the level of baseplate, and there were chemical gradients of major binder elements along the building direction. At the upper half of the specimen, there were relatively less variation in chemical composition and more
EP
homogeneously distributed phases including WC, W2C, η-carbide and FCC metal binder. The hardness of the lower half specimen varied from 711.7 HV1 (bottom of
AC C
the specimen) to 1178.6 HV1 at 1 mm height. For the upper half of the specimen, hardness values could range from 1306.8 HV1 to 1413.4 HV1 and fracture toughness varied from 9.74 MPa m1/2 to 13.29 MPa m1/2. Keywords: additive manufacturing; selective laser melting; cemented carbide; microstructural characterization
1
ACCEPTED MANUSCRIPT
1. Introduction Cemented carbide consisting of hard tungsten carbide particles and ductile metal binder phase can possess extraordinary hardness, wear-resistance, and toughness. Tools made of cemented carbide have been widely used in industrial applications, such as machining, cutting or drilling [1-3]; however, traditional fabrication route
RI PT
through liquid phase sintering can only make mold with symmetrical and simple geometry [4, 5]. Thus, in order to manufacture cemented carbide tools with complex geometry to meet advanced industrial requirements, scientists and engineers are exploring possibilities to employ additive manufacturing (AM) processes [5-12]. For the fabrication of cemented carbides, there are various AM processes that have been
SC
attempted and they can be classified by the need for post-sintering process. For those techniques requiring post-sintering, including 3D gel-printing [5] and Thermoplastic
M AN U
3D Printing [5, 6], the cemented carbide raw materials are blended with organic binders as a slurry for shape-forming. The slurry is then dispatched by a nozzle to form green parts that need to be sintered to remove organic ingredients as well as to obtain the designed geometries. On the other hand, laser engineered net shaping [7, 8], electron beam melting [9], selective laser sintering/melting [13] are those processes without the need for post-sintering process; thus, the powders are free from organic
TE D
ingredients and can be directly consolidated by laser or electron beam via sinter/melting to form near-net-shape parts. Among all techniques, selective laser melting (SLM) and selective laser sintering (SLS) processes have received the most attention. Previous studies have utilized powders with high fraction of carbide
EP
particles (> 80 wt%) [14-16]; bronze-infiltrated WC-9Co composite had been fabricated by SLS process, although the wear resistance of the as-built mold could be
AC C
comparable to its counterpart made by tool steel, the infiltrated WC-Co composite could not be 100% dense due to the nature of SLS process. To further increase the density, partial melting of WC by laser with higher power had been suggested [14]. Uhlmann E. et al. [15] utilized agglomerated and pre-sintered powders for SLM process under various laser parameters; they unveiled that the relative density of cemented carbide parts could only be improved at the cost of much loss of Co by evaporation. They also concluded a conflict between the optimization of the relative density of SLM cemented carbide parts and the remaining content of Co, which could influence the toughness of the built parts. Domashenkov A. et al. [16] adopted conventional and nano-structured WC-12Co composite powders for SLM process and 2
ACCEPTED MANUSCRIPT investigated the influence of different powders on the microstructure and mechanical properties of the as-built cemented carbide; they failed to achieve 100 % dense part; however, they found that cemented carbide parts made from nano-phase powder could possess higher homogeneity in microstructure. Among the aforementioned studies, Co has been the most popular choice as the binder material due to its outstanding intrinsic
RI PT
properties. Furthermore, its moderate solubility in WC not only results in an excellent wettability between Co and WC, but also makes it easier to retain the carbon content during sintering [17, 18]; however, due to the high cost and high toxicity of Co [19], substitution of Co as the binder metal has been a subject of interest for decades [20-24]. Our previous work used Ni as the binder metal for WC-W2C, and a dense
SC
microstructure could be achieved by SLM process [25]. Recently, multi-element alloy such as high entropy alloy (HEA) has been used as the binder metal to fabricate
M AN U
cemented carbide by liquid phase sintering process, and inhabitation of the growth of WC particles during the sintering process by the complex alloy system has been reported [26, 27]. For laser-based processes such as laser welding, laser cladding and SLM/SLS, the properties of the built can be affected by the inter-diffusion or dilution between the baseplate material and the deposited material [28]. However, the proposed sluggish diffusion of the HEA might hinder interdiffusion in cemented
TE D
carbide systems [29, 30] during the SLM process. Previous work tried to laser-clad several layers of AlCoCrCuFeNi HEA onto a Mg-substrate, with proper process parameters, no serious dilution occurred on the HEA composition, which preserved the designed properties of the binder alloy [31]; however, to the best of the authors’
EP
knowledge, to date there is no reported study on attempting to build cemented carbide with high entropy alloy binder by SLM process. This study aims to examine the use
AC C
of HEA as a binder material for cemented carbide by investigating the effect of SLM process on chemical composition of the built and properties; this study can provide a reference for future development of cemented carbide processed by SLM.
2. Experimental procedure A powder mixture consisting of 80 wt.% of spherical cast tungsten carbide powder (TEKMAT WC-45, TEKNA Advanced Materials Inc., Canada) and 20 wt.% of gas-atomized multi-element NiAlCoCrCuFe high entropy alloy powder was utilized in this study. The microstructure of both powders were given in Fig. 1. The binder alloy NiAlCoCrCuFe utilized Ni, Fe, Co as the major elements (> 19 wt.%) 3
ACCEPTED MANUSCRIPT and Al, Cr, Cu as the minor elements (< 12 wt.%). The addition of Ni and Fe aimed to decrease the use of Co; Cr, Al and Cu were for solid solutions and oxidation resistance [27, 32]. The nominal compositions of both forms of powders are given in Table 1, and the phase constitutions of both powders were checked by XRD as given in Fig. 2. The NiAlCoCrCuFe alloy powder comprised a single FCC phase.
RI PT
cast tungsten carbide powder consisted of two phases, i.e. WC and W2C, and the The mean particle sizes of carbide powder and NiAlCoCrCuFe alloy powder were measured by Laser Diffraction Particle Size Analyzer (Coulter LS230, Beckman Coulter Inc., U.S.) as 27.69 µm and 47.88 µm, respectively. The Carr’s
SC
compressibility index of the powder mixture was 3.3 %, which indicated an excellent flowability for powders and its suitability for powder-bed type SLM process.
M AN U
The SLM machine utilized in this experiment was a custom-made setup equipped with an ytterbium fiber laser (λ = 1070, YLR-500-SM-AC, IPG Photonics Co., U.S.). The diameter of the laser spot was 57 µm on the fixed focal plane. The oxygen level of the working chamber was kept below 1500 ppm with a protective argon circulation system to prevent oxidation. A baseplate made of Invar alloy (36 wt.% Ni, 64 wt.% Fe) was used due to its low value of thermal expansion. Cemented
TE D
carbide specimens were built on top of the baseplate as square blocks (15 mm × 15 mm × 3 mm) by a set of optimized laser parameters (140 W of laser power, 90 mm/s of scan speed, 115 µm of hatch distance, and 40 µm of layer thickness). As-built cemented carbide specimens were sectioned from the baseplate and cut
EP
transversely by wire-cut electrical discharge machining to reveal the cross-sectional microstructures. Samples were mounted, mechanically ground and polished by
AC C
diamond grinding disk and diamond polishing suspensions down to 3 µm, followed by a final 0.05 µm alumina suspension polishing. The as-polished microstructure was then etched by Murakami’s reagent (1 g KOH, 1 g K3[Fe(CN)6], 100 ml H2O) for 12 seconds, therefore the present phases could be identified based on their different reaction rates to Murakami’s reagent [33]. For example, W2C should be rapidly etched and shown as dark-color; on the other hand, WC and η-phase were more resistant to Murakami’s reagent and remained to be lighter in color contrast. Microstructure observations were conducted by BX-51 optical microscope (OM, Olympus Co. Ltd., Japan), SU-8010 field-emission scanning electron microscope (FE-SEM, Hitachi Ltd., Tokyo, Japan) and JSM-IT100 SEM (JEOL Ltd., Tokyo, 4
ACCEPTED MANUSCRIPT Japan). Analyses on chemical composition were performed by JSM-IT100 SEM with its attached INCA x-act energy dispersive x-ray spectrometer (EDS, Oxford Instruments plc, UK). Phases were identified by D2 Phaser X-ray diffractometer (XRD, Bruker Co. Ltd., U.S.) with Cu Kα radiation (λ = 1.54 Å), and the measurement was set to scan
RI PT
with a fixed step size of 0.02º. Samples were ground and polished for accurate phase identification. For EBSD analysis, JEOL SM-09010 Cross-Section Polisher (JEOL, Ltd., Japan) was utilized to prepare ultra-fine sample surfaces. The surface polishing was carried out by argon ion under 6 kV of accelerating voltage for 16 h. Crystallographic information for EBSD characterization was obtained from PDF-4+
SC
2015 database (International Center of Diffraction Data, U.S.).
Vickers hardness test was carried out by model HM-115 micro vickers hardness
M AN U
testing machine (Mitutoyo Co., Japan). The Vickers hardness value was calculated according to HV = 1.8544 × f/d2, where f (kgf) is the load force and d (mm) is the mean diagonal length of the indentation. For the hardness profile, a load force of 1 kgf and a load duration of 12 seconds were set. From the original surface level of the baseplate (z = 0), hardness test was conducted every 0.25 mm along the building direction and 5 impressions were made for an average value. For the indentation
TE D
toughness test, a load force of 10 kgf was applied to form adequate indentation impressions and Palmqvist cracks at the corners. The toughness test was carried out at height level of z = 1 mm, 1.5 mm, 2.0 mm and 2.5 mm. Four indentation impressions were made for an average value. The fracture toughness (KIC) was determined based
EP
on formula (1), which was suggested by Niihara [34], Warren and Matzke [35]: (1)
AC C
KIC = β(HW)1/2 MPa m1/2,
where β is a constant and generally equals to 8.89 × 10-2 for tungsten carbide composite [36]; H is the hardness expressed in MPa; W is the crack resistance expressed as (P/Σl), which is the load force (P) over the total length (Σl) of all Palmqvist cracks originated from the corners of the Vickers hardness impression. Phase diagrams were simulated by CALPHAD-based Thermo-Calc software with TCFE7 database [37]. Phase diagrams were constructed according to the chemical composition measured by SEM-EDS. The temperature range (y-axis) and the carbon fraction (x-axis) were set from 1273 K to 3273 K and 1 wt.% to 3.5 wt.%, 5
ACCEPTED MANUSCRIPT
RI PT
respectively to cover the range of interest.
Fig. 1. Microstructure of (a) cast WC/W2C tungsten carbide powder and (b)
SC
gas-atomized NiAlCoCrCuFe alloy powder
M AN U
Table 1. Nominal compositions of WC-45 cast tungsten carbide powder and NiAlCoCrCuFe atomized alloy powder Binder
Ni
Al
Co
Cr
Cu
Fe
wt.%
39.36
4.89
20.29
4.71
11.52
19.23
Element
W
C
Fe
Co
Ni
Others
wt.%
Balance
3.0-4.1
< 0.3
< 0.1
< 0.1
< 0.1
AC C
EP
TE D
WC-45
Element
6
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
Fig. 2. XRD profile of (a) the cast tungsten carbide powder shows a constitution of WC and W2C phase, and (b) the NiAlCoCrCuFe alloy powder consists of a single
TE D
FCC phase.
3. Results and Analysis
3.1. Profile of chemical composition
The as-built SLM cemented carbide and its cross-section micrograph are shown
EP
in Fig. 3 and Fig. 4. The density and the fraction of porosity within were determined by Archimedes’ method as 13.46 g/mm3 and 0.64 %, respectively. Dense SLM
AC C
cemented carbide specimens have been fabricated in this work. The SLM process could induce melting of the baseplate and causing dilution effect on the chemical composition of the cemented carbide. Therefore, the chemical compositions at different heights of the built were analyzed by EDS and listed in Table 2 as well as plotted in Fig. 5 (carbon was excluded due to inaccuracy); about 0.45 mm beneath the baseplate surface possessed the initial Invar alloy composition. At the lower part of the built (from z = 0 mm to z = 1 mm), the weight fraction of Fe and Ni surpassed those of the nominal quantities and compositional gradients of Fe, Ni and W could be identified; however, in the upper part of the specimen (from z = 1 mm to the top surface), the fractions of major binder elements (Fe, Co, Ni) became less than the 7
ACCEPTED MANUSCRIPT nominal composition and remained constant. The measured/nominal ratios of major binder elements decrease from 2.35 (z = 0 mm) to 0.91 (z = 1 mm). Therefore, the lower part of the sample (from z = 0 mm to z = 1 mm) and the upper part of the sample (from z = 1 mm to the top surface) were termed as the region (measured/nominal ratio ≥ 0.91) and the region
(measured/nominal ratio ≤ 0.91),
M AN U
SC
RI PT
respectively, in this study.
AC C
EP
TE D
Fig. 3. As-built SLM cemented carbide specimen
8
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
Fig. 4. Optical micrograph of the cross-section of the as-built cemented carbide and the Vickers hardness impressions (load force = 1 kgf) along the building direction.
9
ACCEPTED MANUSCRIPT Table 2. Chemical compositions along the building direction by EDS analysis (wt.%) Major Alloy Element Site
Height (mm)
Minor Alloy Element
W
Fe
Co
Ni
Cr
Cu
Al
Measured/ Nominal
-
79.41
3.96
4.18
8.10
0.97
2.37
1.01
-
1
-0.45
2.43 ±
56.70 ±
0±
39.29 ±
0±
1.59 ±
0±
-
0.46
1.09
0.53
0.96
0.19
1.24
1.70
6
1.65 2.80
1.00
0.30
0.85
0.29
69.06 ± 10.58 ± 5.01 ± 12.64 ± 0.78 ± 1.19 ± 0.74 ± 1.37
0.66
0.56
0.81
81.18 ±
4.26 ±
3.87 ±
7.65 ±
1.25
0.50
0.57
0.72
85.6 ±
3.64 ±
3.44 ±
5.59 ±
1.13
0.46
0.52
0.64
86.33 ±
3.12 ±
3.44 ±
5.45 ±
1.21
0.50
0.55
0.68
0.31
0.64
0.26
0.67 ± 1.62 ± 0.76 ± 0.35
0.67 0.60 0.65
0.70
0.25
0.57 ± 0.87 ± 0.21 ± 0.36
0.91
0.27
0.62 ± 0.67 ± 0.44 ± 0.33
1.50
0.66
0.27
TE D
5
1.00
0.56
EP
4
0.44
1.08
2.35
AC C
3
0.14
51.66 ± 25.93 ± 2.36 ± 18.04 ± 0.17 ± 0.62 ± 1.21 ±
SC
0
M AN U
2
RI PT
Nom.
Fig. 5. Chemical composition profile of major binder elements and tungsten along the building direction.
10
ACCEPTED MANUSCRIPT 3.2. Microstructure Characterization To analyze the microstructure within region
(z = 0 - 1 mm), cemented carbide
specimen connected to a thin piece of baseplate was prepared; the cross-section micrograph is shown in Fig. 6(a). According to the back-scattered electron image shown in Fig. 6(b), deeply-penetrated molten pools (up to 300 µm in depth) beneath
RI PT
the original baseplate surface could be observed. Some un-melted tungsten carbide powders (indicated by red arrows) sank below the baseplate surface due to higher density within the molten pools. The applied high-power laser has led to melting and mixing of elements from carbide powders, alloy powders and baseplate alloy. η-carbides (M2W4C, M3W3C), W2C and FCC binder phase were the primary phases
SC
as characterized by XRD (Fig. 7). Due to the dilution of Fe and Ni from the baseplate, η-carbides formation were promoted [20]. High fraction of η-carbides was also
M AN U
identified by EBSD near the baseplate surface level (z = 0 mm) as shown in Fig. 8(a) and (b). Based on the image analysis results, an average volume fraction of 55.03 %
EP
TE D
of η-carbide and 44.97 % of metal binder phase were present at z = 0 mm.
AC C
Fig. 6. (a) OM and (b) SEM back-scattered micrographs of the region baseplate surface.
11
near the
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
Fig. 7. Four phases including WC, W2C, η-carbide and FCC metal binder could be identified by XRD within the SLM cemented carbide specimen. Phase constitutions , the region
, and the molten pools (top of the specimen)
TE D
varied between the region
AC C
EP
along the building direction.
Fig. 8. (a) SEM image and (b) EBSD phase mapping result identified the presence of η-carbide within the region
near the baseplate surface. (Yellow: M2W4C, Green:
M3W3C, Red: FCC, Blue: WC, Black: non-indexed point) At region
, microstructure contained faceted WC precipitates embedded in the 12
ACCEPTED MANUSCRIPT matrix, Fig. 9(a), and it shows the matrix featuring a W2C/metal dendritic structure labeled as site 1 in Fig. 9(b). The back-scattered electron image in Fig. 10 shows higher-magnified microstructure, which is at the same region as shown in sites 2 and 3 within Fig. 9(b); except for the light-grey W2C dendrite and the black inter-dendritic metal binder, a dark-grey phase formed heterogeneously and substitute some fractions
RI PT
of metal binder phase. The chemical composition of different phases (sites labeled as A, B, C, D and E) were checked by SEM-EDS and listed in Table 3. Comparing to the XRD analysis results, Fig. 7, site A and C could be identified as W2C phase for its high fraction of tungsten; sites B and D could be characterized as η-carbide for its higher fraction of Fe and Ni elements; site E could be identified as the FCC metal contained 13.14 % of WC,
SC
binder phase. Based on the image analysis results, region
39.61 % of W2C, 19.05 % of FCC binder phase and 28.20 % of η-carbide (at z = 2.0
TE D
M AN U
mm).
Fig. 9. Optical micrographs of the region
microstructure (a) as-polished and (b)
AC C
EP
etched by Murakami’s reagent
Fig. 10. Back-scattered electron image showing SEM micrographs of the region W2C are A and C, η-carbides are B and D, E is the metal binder. 13
.
ACCEPTED MANUSCRIPT Table 3. Chemical composition of different phases as labeled in Fig. 10 (in wt.%) Major Alloy Elements
Minor Alloy Elements
W
Fe
Co
Ni
Cr
Cu
Al
A
96.58 ±
1.06 ±
0.43 ±
1.19 ±
0.66 ±
0.08 ±
0 ± 0.11
0.49
0.19
0.21
0.25
0.17
0.29
82.90 ±
4.45 ±
4.59 ±
6.23 ±
0.85 ±
0.60 ±
0.38 ±
0.57
0.24
0.28
0.32
0.16
0.29
0.12
97.10 ±
1.01 ±
0.53 ±
0.76 ±
0.59 ±
0 ± 0.28
0 ± 0.11
0.39
0.19
0.20
0.24
0.16
84.19 ±
3.61 ±
4.19 ±
6.21 ±
0.56
0.23
0.27
0.32
47.04 ±
9.33 ±
9.08 ±
25.42 ±
0.83
0.30
0.34
0.54
D E
SC
C
0.84 ±
0.68 ±
0.29 ±
0.16
0.29
0.12
0.60 ±
4.81 ±
3.73 ±
0.15
0.37
0.17
M AN U
B
RI PT
Site
The micrograph given in Fig. 11(a) shows the top surface region of the built, which was the last solidified volume of the specimen, featuring keyhole-shaped molten pools at the top of region
. Fig. 11(b) shows the W2C/metal dendritic
TE D
colonies within the molten pools. Un-melted carbide powders were hardly seen in this region, which indicates that adopted high-power laser could fully melt both tungsten carbide powders and alloy powders within the molten pools. The depth of a molten pool could be up to 400 µm, which was 10 times the thickness of a single powder
EP
layer (40 µm) during the spreading process. Fig. 12 shows microstructure at the boundary between the molten pools and the
AC C
inter-pool area. Fine W2C/metal dendritic structure could be identified in the molten pool (left-hand side), and the inter-pool area was heat-affected due to overlapping of laser scanning and showed coarser W2C dendrites and WC precipitates (right-hand side); however, η-carbide was absent in the molten pools and the inter-pool area. Based on the image analysis results, the volume fraction of the dendritic W2C phase and the FCC binder phase were 87.30 % and 12.70 % in average at z = 2.5 mm, respectively.
14
RI PT
ACCEPTED MANUSCRIPT
EP
TE D
M AN U
SC
Fig. 11. (a) as-polished and (b) etched optical micrographs of the molten pool area.
AC C
Fig. 12. Microstructure at the boundary of a molten pool and the inter-pool area. Coarsened WC precipitates and W2C dendrites can be identified in the inter-pool area (right-hand side), while the molten pool remains a fine dendritic structure (left-hand side).
Vickers hardness test was carried out on the cross-section of the specimen along
the building direction, and the test results were drawn as a hardness profile shown in Fig. 13. In the region
, a gradient value of hardness is present and the average
hardness increases from 711.7 HV1 at z = 0 mm to 1178.6 HV1 at z = 1 mm. As for the region
, hardness values vary from 1306.8 HV1 (z = 1.25 mm) to 1413.4 HV1 (z
= 2.50 mm). On the other hand, indentation toughness test was also conducted along 15
ACCEPTED MANUSCRIPT the building direction, and the toughness profile is given in Fig. 13. Microstructure of the indentation impression and the radial cracks formed at the corners are given in Fig. 14. For the region
, its fracture toughness was too high to be determined by
indentation fracture method since no crack formed during the test. The region possesses fracture toughness ranging from 13.29 MPa m1/2 (at z = 1.0 mm) to 9.74
TE D
M AN U
SC
RI PT
MPa m1/2 (at z = 2.5 mm).
Fig. 13. Profile of Vickers hardness and indentation fracture toughness at different
AC C
EP
height level along the building direction.
Fig. 14. Micrographs of Vickers' indentation impressions and induced surface Palmqvist cracks in region II (a) z = 1.0 mm, and (b) z = 2.5 mm (molten pool area). 16
ACCEPTED MANUSCRIPT
4. Discussion According to the composition profiles shown in Table 2 and Fig. 5, the variation of composition within region
and region
can be governed by two mechanisms,
i.e. evaporation of elements during laser scanning and dilution by elements diffused from baseplate Invar alloy, i.e. Ni and Fe. Although the molten pools beneath the top
RI PT
surface of the specimen were not affected by the dilution from the baseplate, deviation from the nominal compositions in this region indicates that high-power laser could cause composition changes due to evaporation of elements, and this phenomena has been reported previously [38]. Within region
(z = 0 – 1 mm) and along the building
direction, the increasing species were W, Co, Cr, Cu, and the decreasing species were
SC
Fe, Ni and Al. Diffusion of Fe and Ni toward the built resulted in an average volume fraction of 55.03 % of η-carbide and 44.97 % of metal binder phase at z = 0 mm. The
M AN U
fractions of Co, Cr and Cu were lower than their nominal fractions at z = 0, possibly due to inter-diffusion toward the baseplate. As the cemented carbide kept building up, the molten pools moved upwards and there was less dilution effect from the baseplate, thus the fractions of other elements, such as W, Co, Cr, Cu were higher than those of the lower regions. In region
(z = 1 – 2.8 mm), W was the only element with
increased contents, and all the other binder elements, e.g. Fe, Co, Ni, Cr, Cu, Al,
TE D
decreased along the building direction, this was caused by the evaporation elements [38, 39], and the increase in W was primarily resulted from the loss of the other elements through vaporization.
The magnitude of elemental contents variation within the region
EP
smaller comparing to those of region
was much
, Fig. 5. Therefore, region II exhibited
relatively more uniform microstructure containing around 13.14 % of WC, 39.61 % of
AC C
W2C, 19.05 % of FCC binder phase and 28.20 % of η-carbide. The measured/nominal weight fraction ratios of the metal binder elements derived from Table 2 can be utilized to quantify the loss for each element in Table 4. The loss in weight fraction for elements are different from each other. According to the work done by Khan and Debroy [40], the rates of element loss during laser welding depend on the vaporization rates of element within the molten pools. The vaporization rates of element can be determined by the following equation [41]: Ji = AcPi/(2πMiRT)1/2
17
(2)
ACCEPTED MANUSCRIPT where Ji (kg-mole/m2 s) is the vaporization rate, Ac is a dimensionless parameter, Pi (N/m2) is the partial pressure of element i, Mi (kg/kg-mole) is the molecular weight of element i, R (kg m2/ s2 kg-mole K) is the gas constant, and T (K) is the temperature. In this study, the vaporization rates of different elements can be proportional to Pi/(Mi)1/2. According to He X. et al.[42], the peak temperature of the molten pools
RI PT
often exceeds the boiling temperature of elements. And the dendritic W2C present in the molten pools indicates that temperature during SLM should reach higher than the melting point of W2C, i.e. 3003 K [43], therefore, the ratios of vaporization rates for each elements can be derived from literatures and listed in Table 4. The major binder elements, Fe, Co and Ni have similar boiling points and can be compared directly
SC
with each other. Among the three major binder elements, Ni possesses the highest evaporation rate and Co exhibits the least evaporation rate, these analyses agree well
M AN U
with composition measurements in Table 2. For the minor elements, such as Cr, Cu, Al, their vapor pressures cannot be accurately estimated at temperature close to 3200 K due to their relatively lower boiling points and insufficient data; however, it can be reasonably assumed that elemental evaporation loss should be higher for those minor elements having lower boiling points. Therefore, the Al content loss was the most point.
TE D
severe among all the binder elements, and this can be resulted from its lowest boiling On the other hand, for around z=0 in region I, high power laser induced key-hole mode melting and melted not only cemented carbide but also a part of baseplate Invar alloy. As the depth of molten pools (300 µm) was larger than the thickness of one
EP
single powder layer (40 µm), a portion of the baseplate as well as the deposited part could be repeatedly re-melted (for the first several layers) during the SLM process.
AC C
Rapid convection and inter-diffusion of elements driven by the concentration gradients took place within the molten pools induced by high power laser [44]. Thus, comparing to the region
, more pronounced compositional gradients of all binder
elements within region
can be attributed to the dilution effect. As for the carbon
content, due to the difficulty to measure accurate carbon content by SEM-EDS, the variation of carbon fraction within different regions is discussed in the next section with the aid of both microstructure examination and predicted phase diagrams.
18
ACCEPTED MANUSCRIPT
Table 4 Properties, vapor pressure and partial pressure of present elements W
Fe
Co
Ni
Cr
Cu
Al
183.84
55.845
58.933
58.693
51.996
63.546
26.982
Boiling Point [45]
5828 K
3134 K
3200 K
3186 K
Nominal Molar
0.537
0.088
0.088
0.17
1.09
0.79
0.82
0.67
< 1 Pa
1.346 ×
1.013 ×
1.207 ×
5
10 Pa
5
Atomic Weight
2944 K
2835 K
2792 K
0.023
0.046
0.046
0.59
0.37
0.21
1.013 ×
1.327 ×
1.045 ×
10 Pa
5
10 Pa
5
10 Pa
105 Pa
At 3200
[47]
At 2944
[48]
[49]
K (B.P. of
At 3200
K (B.P. of
At 2900
At 2800
Co)
K
Cr)
K
K
0.891 ×
2.052 ×
0.233 ×
0.610 ×
0.481 ×
104 Pa
104 Pa
104 Pa
104 Pa
104 Pa
104 Pa
0.591
0.433
1
0.114
0.297
0.234
d ratio of weight Pi (at 3200 K)
5
10 Pa [46] At 3200 K
Ji/JNi
1.184 ×
-
TE D
Pi = XiPi0
M AN U
fraction 0
SC
Fraction (Xi) Nominal/Measure
RI PT
[45]
There were distinct composition gradients within the region
, so compositions
EP
at three positions were chosen for phase diagram simulations, results are shown in Fig. 15(a), 15(b) and 15(c) corresponding to z = 0 mm, 0.44 mm, and 1.00 mm,
AC C
respectively. By contrast, there was no distinct composition gradients found in region , so the composition of the molten pool area beneath the top surface was used to predict phase diagram of region II, Fig. 15(d). At z = 0 mm, due to the absence of WC, the rapidly cooled molten materials
were likely entered a two-phases domain (liquid + η) as labeled in Fig. 15(a) to form η-carbide at around 1900 K. The remaining molten metal solidified afterwards to form the FCC matrix. At z = 0.44 mm, with less degree of dilution and a lack of WC phase, the solidification process might follow the red arrow shown in Fig. 15(b), where materials went through liquid + W + W2C before reaching the final FCC + η-carbide. At z = 1 mm, the solidifying molten materials entered the “liquid + W2C” domain and 19
ACCEPTED MANUSCRIPT then the “η + WC” domain, Fig. 15(c). For region
, because of the decreased
fraction of metal binder (mainly Fe and Ni), the matrix changed to a W2C/metal dendritic structure as shown in Fig. 9(b). According to the phase diagram in Fig. 15(d), W2C dendrite formed first and followed by rapid solidification of the remaining molten metals as inter-dendritic FCC metal phase (route 1). The W2C/metal dendritic
RI PT
structure was then repeatedly heated by the iteratively laser scanning, which led to precipitation and coarsening of faceted WC from the W2C dendrite. Furthermore, at temperature below 1900 K, η-carbide was more thermodynamically stable than W2C. Phase transformation then took place between the dendritic W2C and the
SC
inter-dendritic metal phase, indicated by the following the reaction (3):
(3)
M AN U
FCC + W2C → WC + η-carbide
There was no distinct microstructural gradient along the building direction within region
, it was observed that phase transformation took place mainly within
the heat-affecting zone around the molten pools. At locations away from the molten pools, phase transformation would be halted due to insufficient thermal energy. For the molten pool area near the top surface, phase transformation took place within the
TE D
“Liquid + W2C” and “Liquid + W2C + WC” domains labeled in Fig. 15(d). The molten pool area featured as-solidified microstructure, which follows the route 3 solidification path in Fig. 15(d); for inter-pool areas, they were slightly heat-affected by the neighboring molten pools and followed the route 2 path. WC phase
EP
precipitated and coarsened from the W2C dendrite and at the inter-pool area. η-carbide was hardly seen within the inter-pool area. According to Fig. 15(d), η-carbide is
AC C
thermodynamically-stable at temperature below 1900 K, therefore, outside temperature range and insufficient thermal energy can be two reasons for the absence of η-carbide within the inter-pool area. Based on the analysis above, the carbon content can be estimated from the phase
diagrams as 1 – 1.2 wt% at z = 0 mm, 1.25 – 1.5 wt% at z = 0.44 mm, 2.3 – 2.6 wt% at z = 1 mm for the region WC) for the region region
and 2.6 – 3.3 wt% (Liquid + W2C → liquid + W2C +
. The decrease of carbon fraction was more pronounced in the
, which could be resulted from the dilution effect as well as diffusion of
carbon into the baseplate by inter-diffusion; however, for the region
, due to the
limited dilution effect by the baseplate elements, the carbon content remained close to 20
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
the nominal content.
21
AC C
EP
TE D
M AN U
SC
RI PT
ACCEPTED MANUSCRIPT
Fig. 15. Thermo-Calc phase diagram simulations for the (a) region (b) region
(at z = 0.44 mm) (c) region
(at z = 0 mm),
(at z = 1.00 mm) (d) region
mm).
22
(at z = 1.65
ACCEPTED MANUSCRIPT The mechanical properties of the SLM cemented carbide vary between different regions within the built and can be related to their microstructures. Comparing the chemical composition profile (Fig. 5) with the hardness profile (Fig. 13), within the region
, the decreasing trend of major binder elements (Fe, Ni) corresponds to the
increase in hardness due to increased fractions of carbides. For region
, the weight
which reflected the hardness value plateau in the region
RI PT
fraction of major binder elements remained at 13-14 wt.% along the building direction, . Since the hardness of
cemented carbide can be proportional to the volume fraction of hard carbide phase [50], with increasing fraction of carbides and decreasing fraction of binder phase could lead to increased hardness as expected. As for the molten pool area beneath the inferior to the average value of the region
SC
top surface, there were only a few WC precipitates, and their fracture toughness was . WC phase possesses better fracture
within region
M AN U
toughness than that of W2C phase [51], thus the evenly-distributed WC precipitates could improve the fracture toughness of this region.
Mechanical properties of commercial cemented carbides [52, 53] and those of the present SLM built are summarized in Table 5 for comparisons.; the binder fractions of the SLM built in Table 5 include 30 wt.% at the middle of the regionⅠ(z = 0.5 mm), 14 wt.% at the region Ⅱ (z = 2.0 mm) and 13 wt.% at the top molten pool
TE D
area (z = 2.5 mm). The properties of the region Ⅰ can be compared with that of the commercial cemented carbide containing 25 wt% Co, and the higher hardness of the SLM built could be attributed to the presence of dendritic W2C, which is harder than that of WC [54, 55]. On the other hand, the region Ⅱ and the top molten pool area can
EP
be comparable to that of commercial cemented carbide containing 12 wt% Co. Although the hardness was slightly higher for the SLM-built parts, the more brittle
AC C
W2C dendrite also yielded lesser fracture toughness [56], especially for the top molten pool area.
Based on the simulated phase diagrams as given in Fig. 15, the formation of
brittle W2C and η-carbide were inevitable in the present system. Full-melting of carbide powders resulted in the W2C dendrite; while carbon-deficiency of the system led to formation of η-carbide. Since W2C could decompose into WC and W at 1300 ℃ [43], post heat treatment may induce WC formation for better toughness. Furthermore, compensation of carbon content may also reduce W2C and η-carbide contents by forming WC [22]. Considering the fabrication sequence of the layer-by-layer SLM process, region 23
ACCEPTED MANUSCRIPT Ⅰ with high fraction of metal binder phase and high fracture toughness can serve as a transition part between the region Ⅱ and the baseplate. The region Ⅱ with constant chemical composition and stable mechanical properties should be used as the main body of the SLM cemented carbide. However, the last-solidified molten pool regions with inferior toughness performance should be removed to enhance the
RI PT
performance. Table 5 Comparison of properties of selected conventionally-sintered cemented carbides (Krupp-Widia, Essen) and the present SLM built [53] Conventionally-Sintered
(wt%) Binder (wt%)
25Co
88
75
87
86
70
12
25
13
14
30
1290
780
1330.2
1395.8
988.5
9.7
11.9
-
Vickers Hardness
SC
Carbides
Top Molten
12Co
Pool
Region Ⅱ
M AN U
Type
Present SLM Cemented Carbide
Region Ⅰ
(kg/mm ) Fracture Hardness 1/2
12.7
14.5
EP
(MPa m )
TE D
3
In summary, SLM process yielded a gradient composition and microstructure
AC C
due to dilution effect from the inter-diffusion between the built and the baseplate. Although the dilution effect could only influence certain height of the built, in this case 1mm, evaporation of elements could also alter the composition of the cemented carbide. Furthermore, laser scanning tracks resulted in molten-pools that produced uneven distribution of carbides. Mechanical properties of the cemented carbides built by SLM were affected by the SLM processing characteristics mentioned above, and the most affected part was region I influenced by dilution effect. Furthermore, the application of HEA composition as the binder in this work could not hinder the uneven growth of carbides during laser scanning overlaps, and it had been very difficult to control the composition of NiAlCoCrCuFe high entropy alloy during the 24
ACCEPTED MANUSCRIPT SLM process due to different element has different diffusion characteristic and evaporation rate. The future design of the binder alloy should reduce the difference in boiling temperatures of the alloying elements, so a better compositional control for the binder might be achieved during SLM process. According to Figure 15, regardless of the variation in binder elements content,
RI PT
whether it was diluted by nickel and iron due to interdiffusion or evaporation due to high power laser scanning, the solidification process all involved in the formation of either brittle η or W2C that could render low fracture toughness of the built. It is possible that a designed scanning strategy can control the solidification path in order to induce more WC formation; prolong post heat treatment might also induce more
SC
WC formation. Future studies will be needed in these areas and are currently in
M AN U
progress.
5. Conclusions
In the present study, 80 wt.% of spherical cast tungsten carbide powder mixed with 20 wt.% of gas-atomized NiAlCoCrCuFe high entropy alloy powder was used to fabricate solid block cemented carbide sample by selective laser melting process. The effect of SLM process on the microstructure and mechanical properties have been 1.
TE D
examined and discussed. The followings are conclusions: On the cross-section of the SLM cemented carbide specimen, chemical composition could vary along the building direction due to elemental dilution from the baseplate and evaporation of elements. It can be characterized into two (from z = 0 to z = 1 mm) and region
EP
regions, region
(from z = 1 mm to 3
mm (top surface)). In region
, which was enriched with Fe and Ni from the baseplate material and
AC C
2.
therefore deviated from the nominal composition significantly, up to 55.03 %
η-carbide were embedded in the metal binder matrix. Region
features
W2C/metal dendritic matrix, evenly-distributed WC precipitates and η-carbide
which evolved from W2C dendrite and inter-dendritic metal binder phase. As for
the last-solidified molten pool area at the top of the region
, up to 87.3 % of
dendritic W2C and no η-carbide could be found, since keyhole-shaped molten pools located just beneath the top surface of the specimen were the least-heat-affected area throughout the specimen. 3.
Region
possesses hardness values vary from 711.7 HV1 (z = 0 mm) to 1178.6 25
ACCEPTED MANUSCRIPT HV1 (z = 1 mm); region
possesses hardness values ranging from 1306.8 HV1
(z = 1.25 mm) to 1413.4 HV1 (z = 2.50 mm). Furthermore, the region possesses moderate toughness ranging from 9.74 MPa m1/2 to 13.29 MPa m1/2 and should be considered as the main body of SLM cemented carbide, however, the top molten pool area should also be removed as well as the region I to obtain
RI PT
part with more desirable mechanical properties.
Acknowledgements
Authors would like to thank Prof. Jien-Wei Yeh and Prof. Su-Jien Lin for the use of hardness testing machine, and Prof. Jeng-Gong Duh for the use of JEOL SM-09010
SC
Cross-Section Polisher. This work was supported by the Ministry of Science and Technology (MOST) [grant numbers: 107-2218-E-007-012, 106-3114-E-007-011 and
M AN U
107-2218-E-007-016]; and the “High Entropy Materials Center” from The Featured Areas Research Center Program within the framework of the Higher Education Sprout Project by the Ministry of Education (MOE) and MOST in Taiwan [grant number: MOST 107-3017-F-007-003].
TE D
Declarations of interest: None
References
[1] L.J. Prakash, Application of fine grained tungsten carbide based cemented
EP
carbides, Int. J. Refract. Met. Hard Mater., 13 (1995) 257-264. [2] H.-O. Andrén, Microstructure development during sintering and heat-treatment of
AC C
cemented carbides and cermets, Mater. Chem. Phys., 67 (2001) 209-213. [3] M.J. Kupczyk, J. Komolka, High durability of cutting insert edges made of nanocrystalline cemented carbides, Int. J. Refract. Met. Hard Mater., 49 (2015) 225-231.
[4] K.J.A. Brookes, 3D-printing style additive manufacturing for commercial hardmetals, Met. Powder Rep., 70 (2015) 137-140. [5] X. Zhang, Z. Guo, C. Chen, W. Yang, Additive manufacturing of WC-20Co components by 3D gel-printing, Int. J. Refract. Met. Hard Mater., 70 (2018) 215-223. [6] J.P. U. Scheithauer, S. Weingarten, E. Schwarzer, A. Vornberger, T. Moritz, A. Michaelis, Droplet-Based Additive Manufacturing of Hard Metal Components by 26
ACCEPTED MANUSCRIPT Thermoplastic 3D Printing (T3DP), J. Ceram. Sci. Technol., 8 (2016) 155-160. [7] V.K. Balla, S. Bose, A. Bandyopadhyay, Microstructure and wear properties of laser deposited WC–12%Co composites, Mater. Sci. Eng., A, 527 (2010) 6677-6682. [8] Y. Xiong, J.E. Smugeresky, L. Ajdelsztajn, J.M. Schoenung, Fabrication of WC– Co cermets by laser engineered net shaping, Mater. Sci. Eng., A, 493 (2008) 261-266.
RI PT
[9] H. Peng, C. Liu, H. Guo, Y. Yuan, S. Gong, H. Xu, Fabrication of WCp/NiBSi metal matrix composite by electron beam melting, Mater. Sci. Eng., A, 666 (2016) 320-323.
[10] C.X. Wang, T. Laoui, J. Bonse, P.J. Kruth, B. Lauwers, L. Froyen, Direct Int. J. Adv. Manuf. Technol., 19 (2002) 351-357.
SC
Selective Laser Sintering of Hard Metal Powders: Experimental Study and Simulation, [11] J.P. Kruth, X. Wang, T. Laoui, L. Froyen, Lasers and materials in selective laser
M AN U
sintering, Assembly Autom., 23 (2003) 357-371.
[12] K. Maeda, T.H.C. Childs, Laser sintering (SLS) of hard metal powders for abrasion resistant coatings, J. Mater. Process. Technol., 149 (2004) 609-615. [13] G. Zong, Y. Wu, N. Tran, I. Lee, D. Bourell, J. Beaman, H. Marcus, Direct selective laser sintering of high temperature materials, in:
1992 International Solid
Freeform Fabrication Symposium, 1992.
TE D
[14] S. Kumar, Manufacturing of WC–Co moulds using SLS machine, J. Mater. Process. Technol., 209 (2009) 3840-3848.
[15] E. Uhlmann, A. Bergmann, W. Gridin, Investigation on Additive Manufacturing of Tungsten Carbide-cobalt by Selective Laser Melting, Proc. CIRP, 35 (2015) 8-15.
EP
[16] A. Domashenkov, A. Borbély, I. Smurov, Structural modifications of WC/Co nanophased and conventional powders processed by selective laser melting, Mater.
AC C
Manuf. Process, 32 (2017) 93-100. [17] G.S. Upadhyaya, Cemented tungsten carbides: production, properties and testing, first ed., William Andrew, New York, 1998. [18] H.E. Exner, Physical and chemical nature of cemented carbides, Int. Mater. Rev., 24 (1979) 149-173.
[19] B. Gries, L.J. Prakash, Hard Materials 1: Acute Inhalation Toxicity by Contact Corrosion-The Case of WC-Co, in:
European Congress and Exhibition on Powder
Metallurgy. European PM Conference Proceedings, The European Powder Metallurgy Association, 2007, pp. 189-196. [20] B. Uhrenius, H. Pastor, E. Pauty, On the composition of Fe-Ni-Co-WC-based 27
ACCEPTED MANUSCRIPT cemented carbides, Int. J. Refract. Met. Hard Mater., 15 (1997) 139-149. [21] G. Gille, J. Bredthauer, B. Gries, B. Mende, W. Heinrich, Advanced and new grades of WC and binder powder – their properties and application, Int. J. Refract. Met. Hard Mater., 18 (2000) 87-102. [22] S. Norgren, J. García, A. Blomqvist, L. Yin, Trends in the P/M hard metal
RI PT
industry, Int. J. Refract. Met. Hard Mater., 48 (2015) 31-45. [23] W.D. Schubert, M. Fugger, B. Wittmann, R. Useldinger, Aspects of sintering of cemented carbides with Fe-based binders, Int. J. Refract. Met. Hard Mater., 49 (2015) 110-123.
[24] M. Tarraste, J. Kübarsepp, K. Juhani, A. Mere, M. Kolnes, M. Viljus, B. Maaten,
SC
Ferritic chromium steel as binder metal for WC cemented carbides, Int. J. Refract. Met. Hard Mater., 73 (2018) 183-191.
M AN U
[25] C.-W. Li, K.-C. Chang, A.-C. Yeh, J.-W. Yeh, S.-J. Lin, Microstructure characterization of cemented carbide fabricated by selective laser melting process, Int. J. Refract. Met. Hard Mater., 75 (2018) 225-233.
[26] C.-S. Chen, C.-C. Yang, H.-Y. Chai, J.-W. Yeh, J.L.H. Chau, Novel cermet material of WC/multi-element alloy, Int. J. Refract. Met. Hard Mater., 43 (2014) 200-204.
TE D
[27] B.L. Ezquerra, T.S. Biurrun, L.L. Cabezas, J.M. Sánchez Moreno, F.I. Lopez, R.M. Pampliega, Sintering of WC hardmetals with Ni-Co-Cr-Ti-Al multi-component alloys, Int. J. Refract. Met. Hard Mater., 77 (2018) 44-53. [28] S. Sun, Y. Durandet, M. Brandt, Parametric investigation of pulsed Nd: YAG
EP
laser cladding of stellite 6 on stainless steel, Surface and Coatings Technology, 194 (2005) 225-231.
AC C
[29] C.-W. Tsai, Y.-L. Chen, M.-H. Tsai, J.-W. Yeh, T.-T. Shun, S.-K. Chen, Deformation and annealing behaviors of high-entropy alloy Al0.5CoCrCuFeNi, J. Alloys Compd., 486 (2009) 427-435. [30] S.G. Ma, J.W. Qiao, Z.H. Wang, H.J. Yang, Y. Zhang, Microstructural features and tensile behaviors of the Al0.5CrCuFeNi2 high-entropy alloys by cold rolling and subsequent annealing, Materials & Design, 88 (2015) 1057-1062. [31] T.M. Yue, H. Xie, X. Lin, H.O. Yang, G.H. Meng, Solidification behaviour in laser cladding of AlCoCrCuFeNi high-entropy alloy on magnesium substrates, J. Alloys Compd., 587 (2014) 588-593. [32] C.-J. Tong, M.-R. Chen, J.-W. Yeh, S.-J. Lin, S.-K. Chen, T.-T. Shun, S.-Y. 28
ACCEPTED MANUSCRIPT Chang, Mechanical performance of the AlxCoCrCuFeNi high-entropy alloy system with multiprincipal elements, Metall. Mater. Trans. A, 36 (2005) 1263-1271. [33] J.R. Davis, ASM specialty handbook: tool materials, ASM international, 1995. [34] K. Niihara, A fracture mechanics analysis of indentation-induced Palmqvist crack in ceramics, J. Mater. Sci. Lett., 2 (1983) 221-223.
RI PT
[35] R. Viswanadham, Science of hard materials, first ed., Springer Science & Business Media, New York, 1983.
[36] D.K. Shetty, I.G. Wright, P.N. Mincer, A.H. Clauer, Indentation fracture of WC-Co cermets, J. Mater. Sci., 20 (1985) 1873-1882.
[37] Thermo-Calc Software TCFE7 Steels/Fe-alloys database.
SC
[38] W.E. King, H.D. Barth, V.M. Castillo, G.F. Gallegos, J.W. Gibbs, D.E. Hahn, C. Kamath, A.M. Rubenchik, Observation of keyhole-mode laser melting in laser
M AN U
powder-bed fusion additive manufacturing, J. Mater. Process. Technol., 214 (2014) 2915-2925.
[39] V.-P. Matilainen, H. Piili, A. Salminen, O. Nyrhilä, Preliminary Investigation of Keyhole Phenomena
during Single
Layer
Fabrication
in
Laser Additive
Manufacturing of Stainless Steel, Phys. Procedia, 78 (2015) 377-387. [40] P.A.A. Khan, T. Debroy, Alloying element vaporization and weld pool (1984) 641-644.
TE D
temperature during laser welding of AlSl 202 stainless steel, Metall. Trans. B, 15 [41] S. Dushman, J. Pierce, Scientific foundations of vacuum technique, Phys. Today, 2 (1949) 27.
EP
[42] X. He, T. DebRoy, P. Fuerschbach, Alloying element vaporization during laser spot welding of stainless steel, J. Phys. D: Appl. Phys., 36 (2003) 3079.
AC C
[43] H.O. Pierson, 6 - Carbides of Group VI: Chromium, Molybdenum, and Tungsten Carbides, in:
Handbook of Refractory Carbides and Nitrides, William Andrew
Publishing, Westwood, NJ, 1996, pp. 100-117. [44] M.M. Collur, A. Paul, T. Debroy, Mechanism of alloying element vaporization during laser welding, Metall. Trans. B, 18 (1987) 733-740. [45] D.R. Lide, CRC handbook of chemistry and physics: a ready-reference book of chemical and physical data, CRC press, 1995. [46] P.D. Desai, Thermodynamic Properties of Iron and Silicon, J. Phys. Chem. Ref. Data, 15 (1986) 967-983. [47] P.D. Desai, Thermodynamic properties of nickel, Int. J. Thermophys., 8 (1987) 29
ACCEPTED MANUSCRIPT 763-780. [48] R. Hultgren, P.D. Desai, Selected thermodynamic values and phase diagrams for copper and some of its binary alloys, (1971). [49] P.D. Desai, Thermodynamic properties of aluminum, Int. J. Thermophys., 8 (1987) 621-638. Mater. Sci. Eng., 33 (1978) 125-133.
RI PT
[50] H.C. Lee, J. Gurland, Hardness and deformation of cemented tungsten carbide, [51] H. Taimatsu, S. Sugiyama, Y. Kodaira, Synthesis of W2C by reactive hot pressing and its mechanical properties, Mater. Trans., JIM, 49 (2008) 1256-1261. [52] V. Sarin, Comprehensive hard materials, Newnes, 2014.
SC
[53] G.S. Upadhyaya, Materials science of cemented carbides — an overview, Materials & Design, 22 (2001) 483-489.
M AN U
[54] M. Khechba, F. Hanini, R. Halimi, Study of structural and mechanical properties of tungsten carbides coatings, Nature & Technology, (2011) 9.
[55] E. Zeiler, S. Schwarz, S. Rosiwal, R. Singer, Structural changes of tungsten heating filaments during CVD of diamond, Materials Science and Engineering: A, 335 (2002) 236-245.
[56] L. Jacobs, M. Hyland, M. De Bonte, Study of the influence of microstructural
TE D
properties on the sliding-wear behavior of HVOF and HVAF sprayed WC-cermet
AC C
EP
coatings, Journal of thermal spray technology, 8 (1999) 125-132.
30
ACCEPTED MANUSCRIPT
Highlights
M AN U
SC
RI PT
Dilution effect, evaporation, and laser scanning tracks could affect the built Hardness and toughness varied due to variation in composition and microstructure
TE D
EP
A complex cemented carbide system was subjected to selective laser melting process The cemented carbide system contains a high entropy alloy Composition and microstructure of the cemented carbide were heterogeneous
AC C