- Email: [email protected]
Study on the sintered characteristics and properties of nanostructured WC–15 wt% (Fe–Ni–Co) and WC–15 wt% Co hard metal alloys
Accepted Manuscript Study on the sintered characteristics and properties of nanostructured WC–15 wt% (Fe–Ni–Co) and WC–15 wt% Co hard metal alloys Shih-Hsien Chang, Professor, Ming-Hung Chang, Kuo-Tsung Huang PII:
S0925-8388(15)30516-8
DOI:
10.1016/j.jallcom.2015.07.119
Reference:
JALCOM 34811
To appear in:
Journal of Alloys and Compounds
Received Date: 17 June 2015 Revised Date:
13 July 2015
Accepted Date: 15 July 2015
Please cite this article as: S.-H. Chang, M.-H. Chang, K.-T. Huang, Study on the sintered characteristics and properties of nanostructured WC–15 wt% (Fe–Ni–Co) and WC–15 wt% Co hard metal alloys, Journal of Alloys and Compounds (2015), doi: 10.1016/j.jallcom.2015.07.119. 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.
The following figure shows the fracture morphology of the WC–(Fe–Ni–Co) and WC-Co
ACCEPTED MANUSCRIPT
specimens by means of high-magnification SEM after the KIC tests. Fig. a shows that numerous binder phases (Fe–Ni–Co) existed in the crack areas, which resisted the penetration and extension of the cracks. Due to the bridging effect of the binder phase, the stress concentration of the crack tip will be resolved through plastic deformation; thus, the cracks did not continue to extend. Once the
RI PT
deformation reaches a critical value, the crack propagation occurs. Meanwhile, the binder phase can link together the two crack faces through the bridging process. Although parts of the cracked areas also showed the bridging effect in the WC-Co specimens, as shown by the arrows (Fig. b), the crack
SC
propagation path was not obviously affected. This result corresponds to the tortuosity phenomenon. Consequently, the bridging process suppressed the crack propagation and resulted in the increase in
TE D
M AN U
tortuosity.
EP
SEM observations of the bridging role of the (a) 1300°C sintered WC–(Fe–Ni–Co), and (b) 1350°C
AC C
sintered WC–Co hard metal alloys after KIC tests.
Study on the sintered characteristics andMANUSCRIPT properties of nanostructured WC–15 wt% ACCEPTED (Fe–Ni–Co) and WC–15 wt% Co hard metal alloys
a
RI PT
Shih-Hsien Chang a,*, Ming-Hung Chang a and Kuo-Tsung Huang b
Department of Materials and Mineral Resources Engineering, National Taipei University of
Department of Auto-Mechanics, National Kangshan Agricultural Industrial Senior High
School, Kaohsiung 82049, Taiwan, ROC
TE D
*Corresponding author,
M AN U
b
SC
Technology, Taipei 10608, Taiwan, ROC
Professor
Department of Materials and Mineral Resources Engineering, National Taipei University of
AC C
EP
Technology, 1, Sec. 3, Zhongxiao E. Rd., Taipei 10608, Taiwan, ROC
Communication:
Department of Materials and Mineral Resources Engineering (MMRE2766), National Taipei University of Technology, 1, Sec. 3, Zhongxiao E. Rd., Taipei 10608, Taiwan, ROC TEL: 886-2-27712171 ext 2766 FAX: 886-2-27317185 E-mail: [email protected]
ACCEPTED MANUSCRIPT
Abstract
In this work, four different vacuum sintering temperatures (1250°C, 1300°C, 1350°C and 1400°C) were studied to determine the optimal process parameters of nano WC–15 wt% (Fe–Ni–Co) and WC–15 wt% Co sintered hard metal alloys. Experimental results showed that the optimal
RI PT
sintering temperatures for nano WC–(Fe–Ni–Co) and WC–Co alloys were 1300°C and 1350°C for 1 h, respectively. The sintered nano WC–(Fe–Ni–Co) and WC–Co hard metal alloys showed a good
SC
contiguity of 0.44 and 0.42; hardness was enhanced to HRA 90.83 and 90.92; the transverse rupture
M AN U
strength (TRS) increased to 2567.97 and 2860.08 MPa; and KIC was 16.23 and 12.33 MPa
m,
respectively. Although the nano WC–(Fe–Ni–Co) alloys possessed a slightly lower TRS value, they exhibited superior fracture toughness (KIC) and hardness similar to that of the nano WC–Co material. Significantly, nano WC–(Fe–Ni–Co) alloys could be sintered at a lower temperature and
TE D
still retained their excellent mechanical properties.
AC C
EP
Key words: sintering, nano WC–(Fe–Ni–Co), nano WC–Co, contiguity, TRS and KIC
1
ACCEPTED MANUSCRIPT
1. Introduction
WC–Co hard metal alloys are widely used for machining, cutting, drilling, forming tools and wear resistant parts, due to their exceptional high hardness, excellent wear resistance and better toughness than that of other hard materials. The properties of WC hard metal alloys depend
RI PT
primarily on binder content (such as Co, Fe and Ni) and grain sizes of WC. When the grain size is reduced to a range of submicron meter or nanometer, the hardness and the strength of WC–Co hard
SC
metal alloys increase remarkably, and the toughness improves greatly as well, thus showing an
M AN U
excellent mechanical property of the hard materials [1,2]. For the sake of above reasons, we use the nano-crystalline WC powders for producing nano WC–Co materials with finer microstructures which represent one of the most active fields of research in the hard metal industry, because the
strength and toughness [3-5].
TE D
nanostructured WC–Co alloys possess excellent mechanical properties such as high hardness,
The grain sizes of WC and composition of the binder are the two major factors, which are
EP
responsible for the mechanical properties of the composite. Typical of the WC grain sizes ranges
AC C
from submicron levels to a few microns and the WC–Co alloys contain 3–30 wt% cobalt [6,7]. For the WC hard metal, cobalt is considered to be the optimal binder metal for most applications, although using cobalt as the binder has several disadvantages, including its market-price fluctuations and environmental toxicity. However, since Fe and Ni belong to the same group as Co and can be strengthened by heat treatment, they are considered to be an ideal alternative binder [8,9]. Moreover, it has a great improvement in toughness of the tungsten carbide when the binder phase is Fe–Ni–Co. The substitution of part or all of the cobalt for nickel or nickel and iron has been 2
investigated in recent years in anACCEPTED attempt mainlyMANUSCRIPT to improve the properties of the binder and at the same time to reduce costs of cobalt powder [8,10]. Powder metallurgy (P/M) is a good method for fabrication of high melting material with better mechanical properties. Sintering is a very important procedure to achieve the dense and ultra-fine
RI PT
grain structure by densifying the powder materials. It is well known that is one of the P/M methods, in order to get the high density of the sintered WC–Co hard metal alloys, high sintering
SC
temperatures, normally located at liquid-sintering state, are usually used [8,10]. In this study, nano
M AN U
WC–Co and WC–(Fe–Ni–Co) hard metal alloys were prepared in an innovative way by means of the vacuum sintering process.
In this study, nano WC–Co and WC–(Fe–Ni–Co) hard metal alloys were prepared in an innovative way by means of the vacuum sintering process. Nanostructured WC–15 wt%
TE D
(Fe–Ni–Co) and WC–15 wt% Co hard metal alloys were produced via vacuum sintering using the P/M technique. The effects of the binders and grain sizes on the properties were also of significant
EP
interest. In this research a series of experimental tests were carried out in order to explore the
AC C
microstructural features and mechanical properties of various sintering temperatures on tungsten carbide hard metal alloys. In addition, the feasibility of the commercial manufacturing of nano WC–(Fe–Ni–Co) hard metal alloys through vacuum sintering was evaluated.
3
ACCEPTED MANUSCRIPT
2. Experimental Procedure
In this research, the binder contents of Fe–Ni–Co and Co were 15 wt%, respectively. Various sintering temperatures were examined in order to find the optimal parameters of nanostructured WC–15 wt% (Fe–Ni–Co) and WC–15 wt% Co alloys and to compare the different properties of the
RI PT
two binders (Fe–Ni–Co and Co) in tungsten carbide materials. The binder Fe–Ni–Co contained 11.62 wt% Fe, 3.02 wt% Ni and 0.36 wt% Co, while the binder Co was 15 wt% Co. Meanwhile, the
SC
nano-sized WC–15 wt% (Fe–Ni–Co) and WC–15 wt% Co alloys were designated as
M AN U
WC–(Fe–Ni–Co) and WC–Co, hereafter.
Both WC–(Fe–Ni–Co) and WC–Co composite powders showed an irregularly shaped surface. The powders also displayed a significant agglomeration phenomenon, as shown in Fig. 1. By using the linear intercept method (ASTM E112-12), the mean particle sizes were 500±52 nm and 490±35
TE D
nm, respectively. During the forming process, the WC–(Fe–Ni–Co) and WC–Co alloy powders were put into an alloy steel mold (6 × 6 × 40 mm3) and a vertical force from a hydraulic press was
EP
applied to the mold at a pressure rate of 3.33 MPa/sec, which was maintained at 235.3 MPa for 5
AC C
min. The powders then underwent a sintering process in which the sintering temperatures were 1250°C, 1300°C, 1350°C and 1400°C. The vacuum was maintained at 1.33 × 10-3 Pa and the soaking time was 60 min.
To evaluate the microstructure and mechanical properties of the WC–(Fe–Ni–Co) and WC–Co hard metal alloys via different sintering temperature, the porosity, hardness, transverse rupture strength (TRS) tests, contiguity, fracture toughness KIC and microstructure inspections were performed. Microstructural features of the specimens were examined by optical microscopy (Nikon 4
Eclipse Lv150) and scanning electron microscopy (Hitachi-S4700). Porosity tests followed the ACCEPTED MANUSCRIPT ASTM B311-08 and C830 standards. The hardness of the specimens was measured by Rockwell indenter (HRA, Indentec 8150LK) with loading of 588.4 N, which complied with the CNS 2114 Z8003 standard methods.
RI PT
The Hung Ta universal material test machine (HT-9501A) with a maximum load of 245 kN was used for the transverse rupture strength (TRS) tests (ASTM B528-05). Meanwhile, Rbm was the
SC
TRS, which is determined as the fracture stress in the surface zone. F was maximum fracture load,
M AN U
L was 30 mm, k was chamfer correction factor (normally 1.00-1.02), b and h were 5 mm in the equation Rbm = 3FLk/2bh2, respectively. The specimen dimensions of the TRS test were 5 × 5 × 40 mm3. Moreover, it needs to slightly grind the surface of the specimen and tests at least three pieces. The toughness of the specimens can be also expressed in the term of the crack resistance, which is
TE D
obtained by the applied load (294.3 N) through the micro-hardness tester (VMT-XT) and is calculated the sum of the dividing crack lengths. The fracture toughness KIC can be calculated by
EP
the following equation [10]:
AC C
KIC = 0.15√(HV30/Σl)
where the HV30 is the hardness (N/mm2), Σl is the sum of crack lengths (mm).
5
ACCEPTED MANUSCRIPT
3. Results and Discussion
Fig. 2 shows the XRD patterns of the WC–(Fe–Ni–Co) and WC–Co alloys after the sintering process at various temperatures. There were no obvious differences between the two specimens; only the characteristic peaks of WC and the binder phases (Fe-Ni and Co) could be found, as shown
RI PT
in Figs. 2a and 2b. A previous study [14] pointed out that the specimens are able to obtain the graphite and η phase that result from escaping and accumulating carbides during the
SC
high-temperature sintering process. In this study, there were not high temperatures enough to
M AN U
generate any impurity or η phases in the sintering process; therefore, the influence of η phases could be excluded from the follow-up study and tests. Fig. 2 also shows that the tungsten carbides were dissolved in the binder phase during the sintering process, which caused the diffraction pattern of the binder to offset and shift slightly by a small angle.
TE D
Fig. 3 shows the volume shrinkage ratio and apparent porosity of the WC–(Fe–Ni–Co) and WC–Co alloys after sintering at various temperatures. Fig. 3a shows that the volume shrinkage ratio
EP
of the WC–(Fe–Ni–Co) specimens showed only a slight variation, reaching 46.04% after sintering
AC C
at 1250°C. Increasing the sintering temperature did not significantly increase the volume shrinkage ratio. Normally, a high volume shrinkage ratio of the specimens can be attributed to the liquid-phase generated by liquid-phase sintering (LPS). Thus, the tungsten carbide particles were close to each other and formed a dense sintered body which led to significant shrinkage in the volume. Since iron was the main binder phase of the WC–(Fe–Ni–Co) alloys, the specimens more easily went into the liquid phase under a lower sintering temperature. As the liquid phase is generated, the shrinkage rate accelerates. Fig. 3a also shows that the apparent porosity of the 6
WC–(Fe–Ni–Co) alloys varied slightly after sintering at various temperatures. The specimens ACCEPTED MANUSCRIPT exhibited apparent porosity of 0.04% after sintering at 1250°C. When the sintering temperature was increased to 1300°C, the apparent porosity had a minimum value (0.03%). As a result, it was
at 1300°C, at which the liquid-phase sintering was achieved.
RI PT
reasonable to suggest that the WC–(Fe–Ni–Co) alloys possess better sintering results after sintering
Fig 3b shows the volume shrinkage ratio and apparent porosity of the WC–Co alloys. It was
SC
found that with cobalt as the binder phase, the sintering temperature must be increased above
M AN U
1250°C to generate LPS results. When the sintered specimen was still in the solid-phase sintering (SPS) process, the relative movement of the particles was hard. This was accompanied by greater porosities; thus, the volume shrinkage was lower (30.11%) and the apparent porosity higher (3.97%) after sintering at 1250°C for 1 h. When the sintering temperature was raised to 1350°C, more and
TE D
more liquid binder phases existed between the tungsten carbide particles, resulting in the redistribution of the tungsten carbide and a denser arrangement. Therefore, the higher volume
EP
shrinkage of the sintered WC–Co specimen was achieved. In addition, Fig. 3b also shows that the
AC C
apparent porosity significantly decreased (3.97% → 1.32% → 0.03% → 0.03%) as the sintering temperature increased. As a result, the volume shrinkage and porosity trends were consistent in the WC–Co alloys.
Fig. 4 shows the BEI (Backscattering Electron Image) morphology of the WC–(Fe–Ni–Co) hard metal alloys after sintering at various temperatures. Fig. 4a shows that the tungsten carbide morphology gradually developed into the ideal prism shape after 1250°C sintering. Although the LPS was generated during the 1250°C sintering, a large number of pores still existed between the 7
powder particles. Moreover, it was still possibleMANUSCRIPT to find parts of the binder phases which had not ACCEPTED been uniformly distributed between the tungsten carbides. By increasing the sintering temperature to 1300°C, higher heat energy was generated which provided a better driving force for the binder phases. Thus, the dissolution and re-precipitation effects were effective in improving the uniform
RI PT
distribution of the binder phases, which resulted in the densified microstructure, as shown in Fig. 4b. As the temperature was further increased (1350°C → 1400°C), the excessive grain growth and
SC
coarsening effects resulted in a slight increase in porosity (Figs. 4c and 4d), which was
M AN U
disadvantageous to the mechanical properties.
Fig. 5 shows the BEI morphology observations of the WC–Co specimens after the sintering treatments at various temperatures. The non-uniform distribution of the binder phases and the broken shape of the tungsten carbides were readily apparent in the 1250°C sintered specimens, as
TE D
shown in Fig. 5a. When the sintering temperature was increased to 1300°C, the LPS effect was gradually produced, and the high porosities and coarsening tungsten carbides resulting from the
EP
more refined tungsten carbide dissolution and re-precipitation were displayed, as shown in Fig. 5b.
AC C
Moreover, when the sintering temperature was increased to 1350°C, the pores between the grains were easily filled in the liquid binding phase and enough energy was provided for the particles to move and diffuse. Consequently, the dissolution and re-precipitation effects resulted in the reduced porosity and the optimal densification was obtained, as shown in Fig. 5c. On the other hand, the grain growth and coarsening phenomenon appeared in the 1400°C-sintered WC–Co hard metal alloys, as shown in Fig. 5d. Contiguity is defined as the average fraction of the surface area shared by a carbide particle 8
grain with all neighboring grainsACCEPTED of the same phase. Therefore, by definition, the contiguity can MANUSCRIPT vary between 0 and 1. In this study, the contiguity can be calculated by the following equation [15,16]: Contiguity = 2NWC–WC / (2NWC–WC + 2NWC–Binder)
RI PT
where NWC–WC is the number of intercepts between a random line of unit length and the WC–WC boundaries; and NWC–Binder is the number of intercepts between the same line and the NWC–Binder
SC
interfaces.
M AN U
Since the interface area of WC–WC was relatively weak, when external stress was applied to the specimens, the contact area of the tungsten carbides with each other was easily broken; thus, the contiguity could not be too high. Normally, a high contiguity value is not expected, but neither can it be too low. According to the literature, the contiguity of the carbide phase in WC–Co decreases
TE D
with an increasing Co content, but it does not depend on the size of the WC grains [17]. In this study, we utilized the same WC and binder powders. As a result, the effects on the contiguity could
EP
be only ascribed to the differences in the sintering temperature.
AC C
Table 1 lists the contiguity of the WC–(Fe–Ni–Co) and WC–Co hard metal alloys, which showed a continuously downward trend as the sintering temperature increased. The result can be further compared with Figs. 4 and 5. Fig. 5a shows the non-uniform distribution of the binder phases and the broken shape of the tungsten carbide resulting from the relatively low sintering temperature (1250°C). The contiguity measurements showed wide scatter, and are listed in Table 1. When the sintering temperature was increased (1300°C → 1350°C →1400°C), the contiguity of the WC–Co alloys rapidly decreased (0.59 → 0.42 → 0.28). Fig. 5 also shows that the WC-Co alloys 9
sintered at 1250°C and 1300°C did not completely reach the LPS condition. As a consequence, the ACCEPTED MANUSCRIPT binder phase resulted in the uniform distribution of the WC particles, making it one of the main factors for the decrease in the contiguity. In addition, the LPS of the WC–(Fe–Ni–Co) alloys was produced at lower sintering temperatures (Fig. 4), which should have affected the contiguity.
RI PT
However, the contiguity of the WC–(Fe–Ni–Co) alloys displayed a continuous decrease as the sintering temperature increased, most notably from 1300°C → 1350°C. Thus, a high temperature
SC
could be another factor in decreasing the contiguity. Table 1 also shows that the grain size of the
M AN U
WC–(Fe–Ni–Co) alloys increased (663 → 717 → 830 → 925 nm) as the sintering temperature increased, and that the contiguity decreased (0.47 → 0.44 → 0.32 → 0.30) simultaneously. In this study, parts of the interface of the carbide particles disappeared during the grain growth and coarsening processes, which resulted in the decrease in NWC-WC. As regards the contiguity equation,
TE D
it can be suggested that the contiguity decreased as the grain size increased. A previous study [17] also indicated that the TRS of WC-Co increased when the contiguity decreased from 1 to ~ 0.5, but
EP
that it decreased when the contiguity decreased from ~ 0.5 to 0. The grain growth and coarsening
AC C
phenomenon which caused the contiguity to decline also affected the mechanical properties. Fig. 6 shows the mechanical tests of the WC–(Fe–Ni–Co) and WC–Co hard metal alloys after sintering at various temperatures. Previous literature [18] has indicated that the solid-solution efficacy for iron content of tungsten carbide was best. The binder phase of iron can be improved and have an excellent solution strengthening effect on the WC–(Fe–Ni–Co) alloys. Thus, the hardness of WC–(Fe–Ni–Co) reached 90.61 HRA after sintering at 1250°C for 1 h. There were slight variations in the hardness as the sintering temperature increased. The maximum hardness of 10
the WC–(Fe–Ni–Co) alloys was ACCEPTED 90.83 HRA, which appeared after sintering at 1300°C for 1 h, as MANUSCRIPT shown in Fig. 6a. With further increases in the temperature (1350°C → 1400°C), the hardness declined (90.34 → 90.21 HRA). It is reasonable to suggest that the grain growth and coarsening phenomenon led to a slight decrease in hardness.
RI PT
Fig. 6a also shows the hardness of the WC–Co hard metal alloys after sintering at various temperatures. When the sintering temperature was increased to 1250°C and 1300°C, the LPS was
SC
not fully obtained and the porosity of the specimens was higher; the hardness was 75.73 and 87.54
M AN U
HRA, respectively. The hardness (90.92 HRA) reached the maximum value after sintering at 1350°C for 1 h; however, when the sintering temperature was increased to 1400°C, the hardness (89.96 HRA) began to decline. The hardness trend was consistent with that of the porosity; thus, it
hardness.
TE D
was concluded that the porosity of the sintered specimens played an important role in affecting their
Fig. 6b shows that both TRS values increased and then decreased as the sintering temperature
EP
increased. The optimal TRS levels of the WC–(Fe–Ni–Co) and WC–Co hard metal alloys were
AC C
2567.97 and 2860.08 MPa after sintering at 1300°C and 1350°C for 1 h, respectively. Owing to the lower sintering temperature (1250°C), the sintering behavior and LPS effects of the tungsten carbide were less complete. Not all specimens can achieve the good densification that is disadvantageous to TRS. When the sintering temperature was gradually increased, the sintering behavior was more effective in eliminating pores. The porosities of the WC–(Fe–Ni–Co) and WC–Co alloys were less than 0.03% after sintering at 1300°C and 1350°C for 1 h, which improved the TRS value. Significantly, 1300°C-sintered WC–(Fe–Ni–Co) and the 1350°C-sintered WC–Co 11
alloys almost reached full densification, effectively reducing the strain points along the rupture ACCEPTED MANUSCRIPT mechanism and increasing the TRS value. Although the grain sizes showed a coarsening phenomenon in the sintered specimens (Table 1), a decreasing trend in TRS value did not seem to agree with the increase in grain size of the WC–(Fe–Ni–Co) alloy (Fig. 6b). It is reasonable to
RI PT
suggest that the effects of the LPS process also resulted in the differences in the mechanical property tests (hardness and TRS). This result also corresponded to the contiguity, as the uniformly
SC
distributed binder phase led to the appropriate decline in the contiguity which was advantageous to
M AN U
the TRS. In addition, the excessive grain growth and the decline in the contiguity after the high-temperature sintering (1400°C) resulted in the decrease in TRS.
The TRS values of the WC–(Fe–Ni–Co) and WC–Co hard metal alloys after the optimal sintering process were compared. The TRS of the WC–(Fe–Ni–Co) specimen was slightly lower
TE D
than that of the WC–Co alloys, as the cobalt composition possessed good comminution, which led to a good distribution in the WC matrix and the excellent TRS. In addition, the WC–(Fe–Ni–Co)
EP
alloys contained 3.02 wt% Ni, which possesses a good plastic property. After a lengthy milling
AC C
process, WC–(Fe–Ni–Co) powders can not obtain uniformity and refinement, thus coarse pores are easily produced during the sintering process, which is one of main factors affecting the TRS. Also, the higher vapor pressure of nickel (ten times that of cobalt) at sintering temperature also causes considerable loss of the nickel binder and it is, therefore, necessary to control the working pressure. The loss of nickel in practice has been reported to be 10 wt% or more at sintering temperatures [19,20]. This means that the volume fraction of Fe–Ni can be significantly decreased during the sintering process, which likely causes the binder phase to produce a non-uniform distribution. 12
The fracture toughness (KIC)ACCEPTED is the ability ofMANUSCRIPT the alloy to resist crack propagation. When a crack rapidly expands and extends, the resistance of the material is an indication of the fracture toughness. Fig. 7 shows the fracture toughness of the WC–(Fe–Ni–Co) and WC–Co hard metal alloys after sintering at various temperatures. The optimal KIC value of 16.23 MPa
m was achieved by the
RI PT
WC–(Fe–Ni–Co) hard metal alloys after sintering at 1300°C for 1 h; the optimal KIC value of the WC–Co specimens (sintered at 1350°C) was 12.33 MPa m. Fig. 7 also shows the fracture
SC
toughness of the WC–(Fe–Ni–Co) and WC–Co specimens, which reached a maximum value and
M AN U
then exhibited a downward trend as the sintering temperature increased. The KIC value of WC–(Fe–Ni–Co) was higher than that of the WC–Co specimens after sintering at various temperatures. Because the main binding phase of the WC–(Fe–Ni–Co) hard metal alloys was iron (11.62 wt%). As compared with other metal-based tungsten carbides, the iron content of tungsten
TE D
carbide can be completely dissolved. As a result, the solid-solution strengthening effect on the iron binder is higher than that of the cobalt binder. The literature also indicated that the weight
EP
percentage of the binder phase over the whole 10 wt%; the Fe–Ni-based tungsten carbide possessed
AC C
the higher fracture toughness [21]. Moreover, the fracture toughness of the WC–(Fe–Ni–Co) was significantly decreased as the sintering temperature increased to 1350°C. In this study, the porosity level and high-temperature coarsening phenomenon of the grain size are believed to be related to low fracture toughness. The excessive grain growth and coarsening effects resulted in a slight increase in porosity (Fig. 4c), which were detrimental to the fracture toughness. Fig. 8 shows the SEM observations of the 1300°C-sintered WC–(Fe–Ni–Co) and 1350°C-sintered WC–Co hard metal alloys after the KIC tests. The fracture behavior and 13
mechanisms of the tungsten carbide were impacted mainly along the grains. Normally, cracks along ACCEPTED MANUSCRIPT the grain boundaries happen in order to expand and absorb the fracture energy through the grain boundaries after KIC tests. Thus, Fe–Ni-based tungsten carbide was more effective in absorbing some of the crack energy which led to the higher fracture toughness. In addition, the crack
RI PT
propagation path of WC–(Fe–Ni–Co) (Fig. 8a) seemed to be more complex and relatively curved, as compared with WC-Co (Fig. 8b). It is safe to say that the Co-based tungsten carbide possessed a
SC
relatively high brittleness, which caused the crack extension of the WC–Co specimen to be
M AN U
relatively straight (Fig. 8b). In order to further examine the crack propagation behavior, Fig. 9 shows the fracture morphology of the WC–(Fe–Ni–Co) and WC-Co specimens by means of high-magnification SEM after the KIC tests. Fig. 9a shows that numerous binder phases (Fe–Ni–Co) existed in the crack areas, which resisted the penetration and extension of the cracks. Due to the
TE D
bridging effect of the binder phase, the stress concentration of the crack tip will be resolved through plastic deformation; thus, the cracks did not continue to extend. Once the deformation reaches a
EP
critical value, the crack propagation occurs [22]. Meanwhile, the binder phase can link together the
AC C
two crack faces through the bridging process. Although parts of the cracked areas also showed the bridging effect in the WC-Co specimens, as shown by the arrows (Fig. 9b), the crack propagation path was not obviously affected. This result corresponds to the tortuosity phenomenon. Consequently, the bridging process suppressed the crack propagation and resulted in the increase in tortuosity.
14
ACCEPTED MANUSCRIPT
4. Conclusions
In this study, the main binders of the WC–(Fe–Ni–Co) hard metal alloys were iron and nickel, and with them a lower sintering temperature (1300°C) was achieved; whereas cobalt, as the chief binding phase of the WC-C hard metal alloys, required a relatively high sintering temperature
RI PT
(1350°C) to achieve a complete LPS result. The optimally sintered WC–(Fe–Ni–Co) and WC–Co hard metal alloys showed the hardness enhanced to HRA 90.83 and 90.92; the TRS increased to
SC
2567.97 and 2860.08 MPa; and the KIC of 16.23 and 12.33 MPa m, respectively.
M AN U
Analyzing the contiguity allowed for a better understanding of the bonding phase dispersion. The contiguity of the WC–(Fe–Ni–Co) and WC–Co hard metal alloys tended to decline as the sintering temperature increased. The optimal contiguity of the WC–(Fe–Ni–Co) and the WC–Co hard metal alloys was 0.44 and 0.42, after sintering at 1300°C and 1350°C for 1 h, respectively.
TE D
According to the experimental results, although the TRS value of the WC–(Fe–Ni–Co) hard metal alloys was slightly lower, but they possessed the better hardness and fracture toughness (KIC).
EP
In addition, the WC–(Fe–Ni–Co) hard metal alloys reached a complete LPS under a relatively low
AC C
sintering temperature. Consequently, the adding of an iron–nickel instead of a cobalt binder for tungsten carbides is preferable.
15
ACCEPTED MANUSCRIPT
Acknowledgments
This research is supported by the ASSAB STEELS TAIWAN CO., LTD. The authors would
AC C
EP
TE D
M AN U
SC
RI PT
like to express their appreciation for Dr. Harvard Chen and Mr. Meng-Yu Liu.
16
ACCEPTED MANUSCRIPT
Reference
[1] W.B. Liu, X.Y. Song, J.X. Zhang, F.X. Yin, G.Z. Zhang, J. Alloys Compd. 458 (2008) 366-371. [2] X. Wang, Z.Z. Fang, H.Y. Sohn, Int. J. Refract. Met. Hard Mater. 26 (2008) 232-241. [3] S. Liu, Z.L. Huang, G. Liu, G.B. Yang, Int. J. Refract. Met. Hard Mater. 24 (2006) 461-464.
RI PT
[4] L.L. Shaw, H. Luo, Y. Zhong, Mater. Sci. Eng. A 537 (2012) 39-48.
[5] I. Konyashin, B. Ries, F. Lachmann, Int. J. Refract. Met. Hard Mater. 28 (2010) 489-497. [6] A. Kumar, K. Singh, O.P. Pandey, Ceram. Int. 37 (2011) 1415-1422.
SC
[7] G.H. Lee, S. Kang, J. Alloys Compd. 419 (2006) 281-289.
[8] S.H. Chang, P.Y. Chang, Mater. Sci. Eng. A 606 (2014) 150-156.
Mater. 12 (1993-1994) 199-206.
M AN U
[9] J.M. Guilemany, I. Sanchiz, B.G. Mellor, N. Llorca, J.R. Miguel, Int. J. Refract. Met. Hard
[10] S.H. Chang, S.L. Chen, J. Alloys Compd. 585 (2014) 407-413.
[11] S.H. Chang, C.W. Lu, J.K. Chen, Int. J. Refract. Met. Hard Mater. 35 (2012) 70-75.
TE D
[12] S.H. Chang, T.P. Tang, K.T. Huang, F.C. Tai, Powder Metall. 54 (2011) 507-512. [13] S.H. Chang, S.H. Chen, K.T. Huang, Mater. Trans. 54 (2013) 1857-1862. [14] H.C. Kim, D.Y. Oh and I.J. Shon, Int. J. Refract. Met. Hard Mat. 22 (2004) 197-203.
EP
[15] S. Luyckx, A. Love, Int. J. Refract. Met. Hard Mat. 24 (2006) 75-79. [16] X. Han, N. Sacks, Y.V. Milman, S. Luyckx, Int. J. Refract. Met. Hard Mat. 27 (2009) 274-281.
AC C
[17] J.L. Chermant, A. Deschanvers and F. Osterstock, Powder Metall. 20 (1977) 63-69. [18] G.S. Upadhyaya, S.K. Bhaumik, Mater. Sci. Eng. A 105-106 (1988) 249-256. [19] C.M. Fernandes, A.M.R. Senos, Int. J. Refract. Met. Hard Mat. 29 (2011) 405-418. [20] T.W. Penrice, J. Mater. Shaping Technol. 5 (1987) 35-39. [21] C.S. Kim, T.R. Massa, G.S. Rohrer, J.R. Miguel, Int. J. Refract. Met. Hard Mat. 24 (2006) 89-100. [22] L.S. Sigl, P.A. Mataga, B.J. Dalgleish, R.M. McMeeking, A.G. Evans, Acta Materialia. 36 (1988) 945-953.
17
Table 1 Comparison of the contiguity and grains size of WC–(Fe–Ni–Co) and (b) WC–Co hard
ACCEPTED MANUSCRIPT
metal alloys after various sintering temperature processes.
Sintering temp. (°C)
Contiguity
Grains size (nm)
1250
0.47
663±40
1300
0.44
717±83
1350
0.32
1400
0.30
1250
－
SC
WC-(Fe, Co, Ni)
RI PT
Specimens
1300 WC-Co
AC C
EP
TE D
1400
18
925±53 －
0.59
－
0.42
683±65
0.28
790±49
M AN U
1350
830±71
Legends
ACCEPTED MANUSCRIPT
Figure 1 SEM photographs of the surface morphology: (a) WC–(Fe–Ni–Co), and (b) WC–Co composite powders, respectively.
Figure 2 XRD patterns of (a) WC–(Fe–Ni–Co), and (b) WC–Co alloys after various sintering
RI PT
temperature processes.
Figure 3 Comparison of the volume shrinkage ratio and apparent porosity of (a) WC–(Fe–Ni–Co),
SC
and (b) WC–Co alloys after various sintering temperature processes.
M AN U
Figure 4 BEI morphology observations of WC–(Fe–Ni–Co) specimens by various sintering temperature treatments: (a) 1250°C, (b) 1300°C, (c) 1350°C, and (d) 1400°C.
Figure 5 BEI morphology observations of WC–Co specimens by various sintering temperature
TE D
treatments: (a) 1250°C, (b) 1300°C, (c) 1350°C, and (d) 1400°C.
Figure 6 Comparison of the mechanical tests of WC–(Fe–Ni–Co) and WC–Co hard metal alloys
EP
after various sintering temperature processes: (a) hardness test, and (b) TRS test.
AC C
Figure 7 Comparison of the fracture toughness of WC–(Fe–Ni–Co) and WC–Co hard metal alloys after various sintering temperature processes.
Figure 8 SEM observations of the (a) 1300°C sintered WC–(Fe–Ni–Co), and (b) 1350°C sintered WC–Co hard metal alloys after KIC tests.
Figure 9 SEM observations of the bridging role of the (a) 1300°C sintered WC–(Fe–Ni–Co), and (b) 1350°C sintered WC–Co hard metal alloys after KIC tests.
19
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
Fig. 1
20
ACCEPTED MANUSCRIPT
Fig. 2
SC
RI PT
(a)
AC C
EP
TE D
M AN U
(b)
The color figure on the Web only 21
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
Fig. 3
The color figure on the Web only 22
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
Fig. 4
23
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
Fig. 5
24
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
Fig. 6
The color figure on the Web only
25
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
Fig. 7
The color figure on the Web only
26
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
Fig. 8
27
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
Fig. 9
28
ACCEPTED MANUSCRIPT
1. The sintered nano WC–Co alloy sintered at 1350°C had the highest hardness (HRA 90.92). 2. The sintered nano WC–(Fe–Ni–Co) alloys showed a good contiguity of 0.44. 3. The optimal nano WC–Co sintered alloy possessed the highest TRS value (2860.08 MPa).
RI PT
4. WC–(Fe–Ni–Co) sintered alloy possessed the highest fracture toughness of KIC (16.23 MPam1/2).
AC C
EP
TE D
M AN U
SC
5. The adding of an iron–nickel instead of a cobalt binder for tungsten carbides is preferable.
S0925-8388(15)30516-8
DOI:
10.1016/j.jallcom.2015.07.119
Reference:
JALCOM 34811
To appear in:
Journal of Alloys and Compounds
Received Date: 17 June 2015 Revised Date:
13 July 2015
Accepted Date: 15 July 2015
Please cite this article as: S.-H. Chang, M.-H. Chang, K.-T. Huang, Study on the sintered characteristics and properties of nanostructured WC–15 wt% (Fe–Ni–Co) and WC–15 wt% Co hard metal alloys, Journal of Alloys and Compounds (2015), doi: 10.1016/j.jallcom.2015.07.119. 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.
The following figure shows the fracture morphology of the WC–(Fe–Ni–Co) and WC-Co
ACCEPTED MANUSCRIPT
specimens by means of high-magnification SEM after the KIC tests. Fig. a shows that numerous binder phases (Fe–Ni–Co) existed in the crack areas, which resisted the penetration and extension of the cracks. Due to the bridging effect of the binder phase, the stress concentration of the crack tip will be resolved through plastic deformation; thus, the cracks did not continue to extend. Once the
RI PT
deformation reaches a critical value, the crack propagation occurs. Meanwhile, the binder phase can link together the two crack faces through the bridging process. Although parts of the cracked areas also showed the bridging effect in the WC-Co specimens, as shown by the arrows (Fig. b), the crack
SC
propagation path was not obviously affected. This result corresponds to the tortuosity phenomenon. Consequently, the bridging process suppressed the crack propagation and resulted in the increase in
TE D
M AN U
tortuosity.
EP
SEM observations of the bridging role of the (a) 1300°C sintered WC–(Fe–Ni–Co), and (b) 1350°C
AC C
sintered WC–Co hard metal alloys after KIC tests.
Study on the sintered characteristics andMANUSCRIPT properties of nanostructured WC–15 wt% ACCEPTED (Fe–Ni–Co) and WC–15 wt% Co hard metal alloys
a
RI PT
Shih-Hsien Chang a,*, Ming-Hung Chang a and Kuo-Tsung Huang b
Department of Materials and Mineral Resources Engineering, National Taipei University of
Department of Auto-Mechanics, National Kangshan Agricultural Industrial Senior High
School, Kaohsiung 82049, Taiwan, ROC
TE D
*Corresponding author,
M AN U
b
SC
Technology, Taipei 10608, Taiwan, ROC
Professor
Department of Materials and Mineral Resources Engineering, National Taipei University of
AC C
EP
Technology, 1, Sec. 3, Zhongxiao E. Rd., Taipei 10608, Taiwan, ROC
Communication:
Department of Materials and Mineral Resources Engineering (MMRE2766), National Taipei University of Technology, 1, Sec. 3, Zhongxiao E. Rd., Taipei 10608, Taiwan, ROC TEL: 886-2-27712171 ext 2766 FAX: 886-2-27317185 E-mail: [email protected]
ACCEPTED MANUSCRIPT
Abstract
In this work, four different vacuum sintering temperatures (1250°C, 1300°C, 1350°C and 1400°C) were studied to determine the optimal process parameters of nano WC–15 wt% (Fe–Ni–Co) and WC–15 wt% Co sintered hard metal alloys. Experimental results showed that the optimal
RI PT
sintering temperatures for nano WC–(Fe–Ni–Co) and WC–Co alloys were 1300°C and 1350°C for 1 h, respectively. The sintered nano WC–(Fe–Ni–Co) and WC–Co hard metal alloys showed a good
SC
contiguity of 0.44 and 0.42; hardness was enhanced to HRA 90.83 and 90.92; the transverse rupture
M AN U
strength (TRS) increased to 2567.97 and 2860.08 MPa; and KIC was 16.23 and 12.33 MPa
m,
respectively. Although the nano WC–(Fe–Ni–Co) alloys possessed a slightly lower TRS value, they exhibited superior fracture toughness (KIC) and hardness similar to that of the nano WC–Co material. Significantly, nano WC–(Fe–Ni–Co) alloys could be sintered at a lower temperature and
TE D
still retained their excellent mechanical properties.
AC C
EP
Key words: sintering, nano WC–(Fe–Ni–Co), nano WC–Co, contiguity, TRS and KIC
1
ACCEPTED MANUSCRIPT
1. Introduction
WC–Co hard metal alloys are widely used for machining, cutting, drilling, forming tools and wear resistant parts, due to their exceptional high hardness, excellent wear resistance and better toughness than that of other hard materials. The properties of WC hard metal alloys depend
RI PT
primarily on binder content (such as Co, Fe and Ni) and grain sizes of WC. When the grain size is reduced to a range of submicron meter or nanometer, the hardness and the strength of WC–Co hard
SC
metal alloys increase remarkably, and the toughness improves greatly as well, thus showing an
M AN U
excellent mechanical property of the hard materials [1,2]. For the sake of above reasons, we use the nano-crystalline WC powders for producing nano WC–Co materials with finer microstructures which represent one of the most active fields of research in the hard metal industry, because the
strength and toughness [3-5].
TE D
nanostructured WC–Co alloys possess excellent mechanical properties such as high hardness,
The grain sizes of WC and composition of the binder are the two major factors, which are
EP
responsible for the mechanical properties of the composite. Typical of the WC grain sizes ranges
AC C
from submicron levels to a few microns and the WC–Co alloys contain 3–30 wt% cobalt [6,7]. For the WC hard metal, cobalt is considered to be the optimal binder metal for most applications, although using cobalt as the binder has several disadvantages, including its market-price fluctuations and environmental toxicity. However, since Fe and Ni belong to the same group as Co and can be strengthened by heat treatment, they are considered to be an ideal alternative binder [8,9]. Moreover, it has a great improvement in toughness of the tungsten carbide when the binder phase is Fe–Ni–Co. The substitution of part or all of the cobalt for nickel or nickel and iron has been 2
investigated in recent years in anACCEPTED attempt mainlyMANUSCRIPT to improve the properties of the binder and at the same time to reduce costs of cobalt powder [8,10]. Powder metallurgy (P/M) is a good method for fabrication of high melting material with better mechanical properties. Sintering is a very important procedure to achieve the dense and ultra-fine
RI PT
grain structure by densifying the powder materials. It is well known that is one of the P/M methods, in order to get the high density of the sintered WC–Co hard metal alloys, high sintering
SC
temperatures, normally located at liquid-sintering state, are usually used [8,10]. In this study, nano
M AN U
WC–Co and WC–(Fe–Ni–Co) hard metal alloys were prepared in an innovative way by means of the vacuum sintering process.
In this study, nano WC–Co and WC–(Fe–Ni–Co) hard metal alloys were prepared in an innovative way by means of the vacuum sintering process. Nanostructured WC–15 wt%
TE D
(Fe–Ni–Co) and WC–15 wt% Co hard metal alloys were produced via vacuum sintering using the P/M technique. The effects of the binders and grain sizes on the properties were also of significant
EP
interest. In this research a series of experimental tests were carried out in order to explore the
AC C
microstructural features and mechanical properties of various sintering temperatures on tungsten carbide hard metal alloys. In addition, the feasibility of the commercial manufacturing of nano WC–(Fe–Ni–Co) hard metal alloys through vacuum sintering was evaluated.
3
ACCEPTED MANUSCRIPT
2. Experimental Procedure
In this research, the binder contents of Fe–Ni–Co and Co were 15 wt%, respectively. Various sintering temperatures were examined in order to find the optimal parameters of nanostructured WC–15 wt% (Fe–Ni–Co) and WC–15 wt% Co alloys and to compare the different properties of the
RI PT
two binders (Fe–Ni–Co and Co) in tungsten carbide materials. The binder Fe–Ni–Co contained 11.62 wt% Fe, 3.02 wt% Ni and 0.36 wt% Co, while the binder Co was 15 wt% Co. Meanwhile, the
SC
nano-sized WC–15 wt% (Fe–Ni–Co) and WC–15 wt% Co alloys were designated as
M AN U
WC–(Fe–Ni–Co) and WC–Co, hereafter.
Both WC–(Fe–Ni–Co) and WC–Co composite powders showed an irregularly shaped surface. The powders also displayed a significant agglomeration phenomenon, as shown in Fig. 1. By using the linear intercept method (ASTM E112-12), the mean particle sizes were 500±52 nm and 490±35
TE D
nm, respectively. During the forming process, the WC–(Fe–Ni–Co) and WC–Co alloy powders were put into an alloy steel mold (6 × 6 × 40 mm3) and a vertical force from a hydraulic press was
EP
applied to the mold at a pressure rate of 3.33 MPa/sec, which was maintained at 235.3 MPa for 5
AC C
min. The powders then underwent a sintering process in which the sintering temperatures were 1250°C, 1300°C, 1350°C and 1400°C. The vacuum was maintained at 1.33 × 10-3 Pa and the soaking time was 60 min.
To evaluate the microstructure and mechanical properties of the WC–(Fe–Ni–Co) and WC–Co hard metal alloys via different sintering temperature, the porosity, hardness, transverse rupture strength (TRS) tests, contiguity, fracture toughness KIC and microstructure inspections were performed. Microstructural features of the specimens were examined by optical microscopy (Nikon 4
Eclipse Lv150) and scanning electron microscopy (Hitachi-S4700). Porosity tests followed the ACCEPTED MANUSCRIPT ASTM B311-08 and C830 standards. The hardness of the specimens was measured by Rockwell indenter (HRA, Indentec 8150LK) with loading of 588.4 N, which complied with the CNS 2114 Z8003 standard methods.
RI PT
The Hung Ta universal material test machine (HT-9501A) with a maximum load of 245 kN was used for the transverse rupture strength (TRS) tests (ASTM B528-05). Meanwhile, Rbm was the
SC
TRS, which is determined as the fracture stress in the surface zone. F was maximum fracture load,
M AN U
L was 30 mm, k was chamfer correction factor (normally 1.00-1.02), b and h were 5 mm in the equation Rbm = 3FLk/2bh2, respectively. The specimen dimensions of the TRS test were 5 × 5 × 40 mm3. Moreover, it needs to slightly grind the surface of the specimen and tests at least three pieces. The toughness of the specimens can be also expressed in the term of the crack resistance, which is
TE D
obtained by the applied load (294.3 N) through the micro-hardness tester (VMT-XT) and is calculated the sum of the dividing crack lengths. The fracture toughness KIC can be calculated by
EP
the following equation [10]:
AC C
KIC = 0.15√(HV30/Σl)
where the HV30 is the hardness (N/mm2), Σl is the sum of crack lengths (mm).
5
ACCEPTED MANUSCRIPT
3. Results and Discussion
Fig. 2 shows the XRD patterns of the WC–(Fe–Ni–Co) and WC–Co alloys after the sintering process at various temperatures. There were no obvious differences between the two specimens; only the characteristic peaks of WC and the binder phases (Fe-Ni and Co) could be found, as shown
RI PT
in Figs. 2a and 2b. A previous study [14] pointed out that the specimens are able to obtain the graphite and η phase that result from escaping and accumulating carbides during the
SC
high-temperature sintering process. In this study, there were not high temperatures enough to
M AN U
generate any impurity or η phases in the sintering process; therefore, the influence of η phases could be excluded from the follow-up study and tests. Fig. 2 also shows that the tungsten carbides were dissolved in the binder phase during the sintering process, which caused the diffraction pattern of the binder to offset and shift slightly by a small angle.
TE D
Fig. 3 shows the volume shrinkage ratio and apparent porosity of the WC–(Fe–Ni–Co) and WC–Co alloys after sintering at various temperatures. Fig. 3a shows that the volume shrinkage ratio
EP
of the WC–(Fe–Ni–Co) specimens showed only a slight variation, reaching 46.04% after sintering
AC C
at 1250°C. Increasing the sintering temperature did not significantly increase the volume shrinkage ratio. Normally, a high volume shrinkage ratio of the specimens can be attributed to the liquid-phase generated by liquid-phase sintering (LPS). Thus, the tungsten carbide particles were close to each other and formed a dense sintered body which led to significant shrinkage in the volume. Since iron was the main binder phase of the WC–(Fe–Ni–Co) alloys, the specimens more easily went into the liquid phase under a lower sintering temperature. As the liquid phase is generated, the shrinkage rate accelerates. Fig. 3a also shows that the apparent porosity of the 6
WC–(Fe–Ni–Co) alloys varied slightly after sintering at various temperatures. The specimens ACCEPTED MANUSCRIPT exhibited apparent porosity of 0.04% after sintering at 1250°C. When the sintering temperature was increased to 1300°C, the apparent porosity had a minimum value (0.03%). As a result, it was
at 1300°C, at which the liquid-phase sintering was achieved.
RI PT
reasonable to suggest that the WC–(Fe–Ni–Co) alloys possess better sintering results after sintering
Fig 3b shows the volume shrinkage ratio and apparent porosity of the WC–Co alloys. It was
SC
found that with cobalt as the binder phase, the sintering temperature must be increased above
M AN U
1250°C to generate LPS results. When the sintered specimen was still in the solid-phase sintering (SPS) process, the relative movement of the particles was hard. This was accompanied by greater porosities; thus, the volume shrinkage was lower (30.11%) and the apparent porosity higher (3.97%) after sintering at 1250°C for 1 h. When the sintering temperature was raised to 1350°C, more and
TE D
more liquid binder phases existed between the tungsten carbide particles, resulting in the redistribution of the tungsten carbide and a denser arrangement. Therefore, the higher volume
EP
shrinkage of the sintered WC–Co specimen was achieved. In addition, Fig. 3b also shows that the
AC C
apparent porosity significantly decreased (3.97% → 1.32% → 0.03% → 0.03%) as the sintering temperature increased. As a result, the volume shrinkage and porosity trends were consistent in the WC–Co alloys.
Fig. 4 shows the BEI (Backscattering Electron Image) morphology of the WC–(Fe–Ni–Co) hard metal alloys after sintering at various temperatures. Fig. 4a shows that the tungsten carbide morphology gradually developed into the ideal prism shape after 1250°C sintering. Although the LPS was generated during the 1250°C sintering, a large number of pores still existed between the 7
powder particles. Moreover, it was still possibleMANUSCRIPT to find parts of the binder phases which had not ACCEPTED been uniformly distributed between the tungsten carbides. By increasing the sintering temperature to 1300°C, higher heat energy was generated which provided a better driving force for the binder phases. Thus, the dissolution and re-precipitation effects were effective in improving the uniform
RI PT
distribution of the binder phases, which resulted in the densified microstructure, as shown in Fig. 4b. As the temperature was further increased (1350°C → 1400°C), the excessive grain growth and
SC
coarsening effects resulted in a slight increase in porosity (Figs. 4c and 4d), which was
M AN U
disadvantageous to the mechanical properties.
Fig. 5 shows the BEI morphology observations of the WC–Co specimens after the sintering treatments at various temperatures. The non-uniform distribution of the binder phases and the broken shape of the tungsten carbides were readily apparent in the 1250°C sintered specimens, as
TE D
shown in Fig. 5a. When the sintering temperature was increased to 1300°C, the LPS effect was gradually produced, and the high porosities and coarsening tungsten carbides resulting from the
EP
more refined tungsten carbide dissolution and re-precipitation were displayed, as shown in Fig. 5b.
AC C
Moreover, when the sintering temperature was increased to 1350°C, the pores between the grains were easily filled in the liquid binding phase and enough energy was provided for the particles to move and diffuse. Consequently, the dissolution and re-precipitation effects resulted in the reduced porosity and the optimal densification was obtained, as shown in Fig. 5c. On the other hand, the grain growth and coarsening phenomenon appeared in the 1400°C-sintered WC–Co hard metal alloys, as shown in Fig. 5d. Contiguity is defined as the average fraction of the surface area shared by a carbide particle 8
grain with all neighboring grainsACCEPTED of the same phase. Therefore, by definition, the contiguity can MANUSCRIPT vary between 0 and 1. In this study, the contiguity can be calculated by the following equation [15,16]: Contiguity = 2NWC–WC / (2NWC–WC + 2NWC–Binder)
RI PT
where NWC–WC is the number of intercepts between a random line of unit length and the WC–WC boundaries; and NWC–Binder is the number of intercepts between the same line and the NWC–Binder
SC
interfaces.
M AN U
Since the interface area of WC–WC was relatively weak, when external stress was applied to the specimens, the contact area of the tungsten carbides with each other was easily broken; thus, the contiguity could not be too high. Normally, a high contiguity value is not expected, but neither can it be too low. According to the literature, the contiguity of the carbide phase in WC–Co decreases
TE D
with an increasing Co content, but it does not depend on the size of the WC grains [17]. In this study, we utilized the same WC and binder powders. As a result, the effects on the contiguity could
EP
be only ascribed to the differences in the sintering temperature.
AC C
Table 1 lists the contiguity of the WC–(Fe–Ni–Co) and WC–Co hard metal alloys, which showed a continuously downward trend as the sintering temperature increased. The result can be further compared with Figs. 4 and 5. Fig. 5a shows the non-uniform distribution of the binder phases and the broken shape of the tungsten carbide resulting from the relatively low sintering temperature (1250°C). The contiguity measurements showed wide scatter, and are listed in Table 1. When the sintering temperature was increased (1300°C → 1350°C →1400°C), the contiguity of the WC–Co alloys rapidly decreased (0.59 → 0.42 → 0.28). Fig. 5 also shows that the WC-Co alloys 9
sintered at 1250°C and 1300°C did not completely reach the LPS condition. As a consequence, the ACCEPTED MANUSCRIPT binder phase resulted in the uniform distribution of the WC particles, making it one of the main factors for the decrease in the contiguity. In addition, the LPS of the WC–(Fe–Ni–Co) alloys was produced at lower sintering temperatures (Fig. 4), which should have affected the contiguity.
RI PT
However, the contiguity of the WC–(Fe–Ni–Co) alloys displayed a continuous decrease as the sintering temperature increased, most notably from 1300°C → 1350°C. Thus, a high temperature
SC
could be another factor in decreasing the contiguity. Table 1 also shows that the grain size of the
M AN U
WC–(Fe–Ni–Co) alloys increased (663 → 717 → 830 → 925 nm) as the sintering temperature increased, and that the contiguity decreased (0.47 → 0.44 → 0.32 → 0.30) simultaneously. In this study, parts of the interface of the carbide particles disappeared during the grain growth and coarsening processes, which resulted in the decrease in NWC-WC. As regards the contiguity equation,
TE D
it can be suggested that the contiguity decreased as the grain size increased. A previous study [17] also indicated that the TRS of WC-Co increased when the contiguity decreased from 1 to ~ 0.5, but
EP
that it decreased when the contiguity decreased from ~ 0.5 to 0. The grain growth and coarsening
AC C
phenomenon which caused the contiguity to decline also affected the mechanical properties. Fig. 6 shows the mechanical tests of the WC–(Fe–Ni–Co) and WC–Co hard metal alloys after sintering at various temperatures. Previous literature [18] has indicated that the solid-solution efficacy for iron content of tungsten carbide was best. The binder phase of iron can be improved and have an excellent solution strengthening effect on the WC–(Fe–Ni–Co) alloys. Thus, the hardness of WC–(Fe–Ni–Co) reached 90.61 HRA after sintering at 1250°C for 1 h. There were slight variations in the hardness as the sintering temperature increased. The maximum hardness of 10
the WC–(Fe–Ni–Co) alloys was ACCEPTED 90.83 HRA, which appeared after sintering at 1300°C for 1 h, as MANUSCRIPT shown in Fig. 6a. With further increases in the temperature (1350°C → 1400°C), the hardness declined (90.34 → 90.21 HRA). It is reasonable to suggest that the grain growth and coarsening phenomenon led to a slight decrease in hardness.
RI PT
Fig. 6a also shows the hardness of the WC–Co hard metal alloys after sintering at various temperatures. When the sintering temperature was increased to 1250°C and 1300°C, the LPS was
SC
not fully obtained and the porosity of the specimens was higher; the hardness was 75.73 and 87.54
M AN U
HRA, respectively. The hardness (90.92 HRA) reached the maximum value after sintering at 1350°C for 1 h; however, when the sintering temperature was increased to 1400°C, the hardness (89.96 HRA) began to decline. The hardness trend was consistent with that of the porosity; thus, it
hardness.
TE D
was concluded that the porosity of the sintered specimens played an important role in affecting their
Fig. 6b shows that both TRS values increased and then decreased as the sintering temperature
EP
increased. The optimal TRS levels of the WC–(Fe–Ni–Co) and WC–Co hard metal alloys were
AC C
2567.97 and 2860.08 MPa after sintering at 1300°C and 1350°C for 1 h, respectively. Owing to the lower sintering temperature (1250°C), the sintering behavior and LPS effects of the tungsten carbide were less complete. Not all specimens can achieve the good densification that is disadvantageous to TRS. When the sintering temperature was gradually increased, the sintering behavior was more effective in eliminating pores. The porosities of the WC–(Fe–Ni–Co) and WC–Co alloys were less than 0.03% after sintering at 1300°C and 1350°C for 1 h, which improved the TRS value. Significantly, 1300°C-sintered WC–(Fe–Ni–Co) and the 1350°C-sintered WC–Co 11
alloys almost reached full densification, effectively reducing the strain points along the rupture ACCEPTED MANUSCRIPT mechanism and increasing the TRS value. Although the grain sizes showed a coarsening phenomenon in the sintered specimens (Table 1), a decreasing trend in TRS value did not seem to agree with the increase in grain size of the WC–(Fe–Ni–Co) alloy (Fig. 6b). It is reasonable to
RI PT
suggest that the effects of the LPS process also resulted in the differences in the mechanical property tests (hardness and TRS). This result also corresponded to the contiguity, as the uniformly
SC
distributed binder phase led to the appropriate decline in the contiguity which was advantageous to
M AN U
the TRS. In addition, the excessive grain growth and the decline in the contiguity after the high-temperature sintering (1400°C) resulted in the decrease in TRS.
The TRS values of the WC–(Fe–Ni–Co) and WC–Co hard metal alloys after the optimal sintering process were compared. The TRS of the WC–(Fe–Ni–Co) specimen was slightly lower
TE D
than that of the WC–Co alloys, as the cobalt composition possessed good comminution, which led to a good distribution in the WC matrix and the excellent TRS. In addition, the WC–(Fe–Ni–Co)
EP
alloys contained 3.02 wt% Ni, which possesses a good plastic property. After a lengthy milling
AC C
process, WC–(Fe–Ni–Co) powders can not obtain uniformity and refinement, thus coarse pores are easily produced during the sintering process, which is one of main factors affecting the TRS. Also, the higher vapor pressure of nickel (ten times that of cobalt) at sintering temperature also causes considerable loss of the nickel binder and it is, therefore, necessary to control the working pressure. The loss of nickel in practice has been reported to be 10 wt% or more at sintering temperatures [19,20]. This means that the volume fraction of Fe–Ni can be significantly decreased during the sintering process, which likely causes the binder phase to produce a non-uniform distribution. 12
The fracture toughness (KIC)ACCEPTED is the ability ofMANUSCRIPT the alloy to resist crack propagation. When a crack rapidly expands and extends, the resistance of the material is an indication of the fracture toughness. Fig. 7 shows the fracture toughness of the WC–(Fe–Ni–Co) and WC–Co hard metal alloys after sintering at various temperatures. The optimal KIC value of 16.23 MPa
m was achieved by the
RI PT
WC–(Fe–Ni–Co) hard metal alloys after sintering at 1300°C for 1 h; the optimal KIC value of the WC–Co specimens (sintered at 1350°C) was 12.33 MPa m. Fig. 7 also shows the fracture
SC
toughness of the WC–(Fe–Ni–Co) and WC–Co specimens, which reached a maximum value and
M AN U
then exhibited a downward trend as the sintering temperature increased. The KIC value of WC–(Fe–Ni–Co) was higher than that of the WC–Co specimens after sintering at various temperatures. Because the main binding phase of the WC–(Fe–Ni–Co) hard metal alloys was iron (11.62 wt%). As compared with other metal-based tungsten carbides, the iron content of tungsten
TE D
carbide can be completely dissolved. As a result, the solid-solution strengthening effect on the iron binder is higher than that of the cobalt binder. The literature also indicated that the weight
EP
percentage of the binder phase over the whole 10 wt%; the Fe–Ni-based tungsten carbide possessed
AC C
the higher fracture toughness [21]. Moreover, the fracture toughness of the WC–(Fe–Ni–Co) was significantly decreased as the sintering temperature increased to 1350°C. In this study, the porosity level and high-temperature coarsening phenomenon of the grain size are believed to be related to low fracture toughness. The excessive grain growth and coarsening effects resulted in a slight increase in porosity (Fig. 4c), which were detrimental to the fracture toughness. Fig. 8 shows the SEM observations of the 1300°C-sintered WC–(Fe–Ni–Co) and 1350°C-sintered WC–Co hard metal alloys after the KIC tests. The fracture behavior and 13
mechanisms of the tungsten carbide were impacted mainly along the grains. Normally, cracks along ACCEPTED MANUSCRIPT the grain boundaries happen in order to expand and absorb the fracture energy through the grain boundaries after KIC tests. Thus, Fe–Ni-based tungsten carbide was more effective in absorbing some of the crack energy which led to the higher fracture toughness. In addition, the crack
RI PT
propagation path of WC–(Fe–Ni–Co) (Fig. 8a) seemed to be more complex and relatively curved, as compared with WC-Co (Fig. 8b). It is safe to say that the Co-based tungsten carbide possessed a
SC
relatively high brittleness, which caused the crack extension of the WC–Co specimen to be
M AN U
relatively straight (Fig. 8b). In order to further examine the crack propagation behavior, Fig. 9 shows the fracture morphology of the WC–(Fe–Ni–Co) and WC-Co specimens by means of high-magnification SEM after the KIC tests. Fig. 9a shows that numerous binder phases (Fe–Ni–Co) existed in the crack areas, which resisted the penetration and extension of the cracks. Due to the
TE D
bridging effect of the binder phase, the stress concentration of the crack tip will be resolved through plastic deformation; thus, the cracks did not continue to extend. Once the deformation reaches a
EP
critical value, the crack propagation occurs [22]. Meanwhile, the binder phase can link together the
AC C
two crack faces through the bridging process. Although parts of the cracked areas also showed the bridging effect in the WC-Co specimens, as shown by the arrows (Fig. 9b), the crack propagation path was not obviously affected. This result corresponds to the tortuosity phenomenon. Consequently, the bridging process suppressed the crack propagation and resulted in the increase in tortuosity.
14
ACCEPTED MANUSCRIPT
4. Conclusions
In this study, the main binders of the WC–(Fe–Ni–Co) hard metal alloys were iron and nickel, and with them a lower sintering temperature (1300°C) was achieved; whereas cobalt, as the chief binding phase of the WC-C hard metal alloys, required a relatively high sintering temperature
RI PT
(1350°C) to achieve a complete LPS result. The optimally sintered WC–(Fe–Ni–Co) and WC–Co hard metal alloys showed the hardness enhanced to HRA 90.83 and 90.92; the TRS increased to
SC
2567.97 and 2860.08 MPa; and the KIC of 16.23 and 12.33 MPa m, respectively.
M AN U
Analyzing the contiguity allowed for a better understanding of the bonding phase dispersion. The contiguity of the WC–(Fe–Ni–Co) and WC–Co hard metal alloys tended to decline as the sintering temperature increased. The optimal contiguity of the WC–(Fe–Ni–Co) and the WC–Co hard metal alloys was 0.44 and 0.42, after sintering at 1300°C and 1350°C for 1 h, respectively.
TE D
According to the experimental results, although the TRS value of the WC–(Fe–Ni–Co) hard metal alloys was slightly lower, but they possessed the better hardness and fracture toughness (KIC).
EP
In addition, the WC–(Fe–Ni–Co) hard metal alloys reached a complete LPS under a relatively low
AC C
sintering temperature. Consequently, the adding of an iron–nickel instead of a cobalt binder for tungsten carbides is preferable.
15
ACCEPTED MANUSCRIPT
Acknowledgments
This research is supported by the ASSAB STEELS TAIWAN CO., LTD. The authors would
AC C
EP
TE D
M AN U
SC
RI PT
like to express their appreciation for Dr. Harvard Chen and Mr. Meng-Yu Liu.
16
ACCEPTED MANUSCRIPT
Reference
[1] W.B. Liu, X.Y. Song, J.X. Zhang, F.X. Yin, G.Z. Zhang, J. Alloys Compd. 458 (2008) 366-371. [2] X. Wang, Z.Z. Fang, H.Y. Sohn, Int. J. Refract. Met. Hard Mater. 26 (2008) 232-241. [3] S. Liu, Z.L. Huang, G. Liu, G.B. Yang, Int. J. Refract. Met. Hard Mater. 24 (2006) 461-464.
RI PT
[4] L.L. Shaw, H. Luo, Y. Zhong, Mater. Sci. Eng. A 537 (2012) 39-48.
[5] I. Konyashin, B. Ries, F. Lachmann, Int. J. Refract. Met. Hard Mater. 28 (2010) 489-497. [6] A. Kumar, K. Singh, O.P. Pandey, Ceram. Int. 37 (2011) 1415-1422.
SC
[7] G.H. Lee, S. Kang, J. Alloys Compd. 419 (2006) 281-289.
[8] S.H. Chang, P.Y. Chang, Mater. Sci. Eng. A 606 (2014) 150-156.
Mater. 12 (1993-1994) 199-206.
M AN U
[9] J.M. Guilemany, I. Sanchiz, B.G. Mellor, N. Llorca, J.R. Miguel, Int. J. Refract. Met. Hard
[10] S.H. Chang, S.L. Chen, J. Alloys Compd. 585 (2014) 407-413.
[11] S.H. Chang, C.W. Lu, J.K. Chen, Int. J. Refract. Met. Hard Mater. 35 (2012) 70-75.
TE D
[12] S.H. Chang, T.P. Tang, K.T. Huang, F.C. Tai, Powder Metall. 54 (2011) 507-512. [13] S.H. Chang, S.H. Chen, K.T. Huang, Mater. Trans. 54 (2013) 1857-1862. [14] H.C. Kim, D.Y. Oh and I.J. Shon, Int. J. Refract. Met. Hard Mat. 22 (2004) 197-203.
EP
[15] S. Luyckx, A. Love, Int. J. Refract. Met. Hard Mat. 24 (2006) 75-79. [16] X. Han, N. Sacks, Y.V. Milman, S. Luyckx, Int. J. Refract. Met. Hard Mat. 27 (2009) 274-281.
AC C
[17] J.L. Chermant, A. Deschanvers and F. Osterstock, Powder Metall. 20 (1977) 63-69. [18] G.S. Upadhyaya, S.K. Bhaumik, Mater. Sci. Eng. A 105-106 (1988) 249-256. [19] C.M. Fernandes, A.M.R. Senos, Int. J. Refract. Met. Hard Mat. 29 (2011) 405-418. [20] T.W. Penrice, J. Mater. Shaping Technol. 5 (1987) 35-39. [21] C.S. Kim, T.R. Massa, G.S. Rohrer, J.R. Miguel, Int. J. Refract. Met. Hard Mat. 24 (2006) 89-100. [22] L.S. Sigl, P.A. Mataga, B.J. Dalgleish, R.M. McMeeking, A.G. Evans, Acta Materialia. 36 (1988) 945-953.
17
Table 1 Comparison of the contiguity and grains size of WC–(Fe–Ni–Co) and (b) WC–Co hard
ACCEPTED MANUSCRIPT
metal alloys after various sintering temperature processes.
Sintering temp. (°C)
Contiguity
Grains size (nm)
1250
0.47
663±40
1300
0.44
717±83
1350
0.32
1400
0.30
1250
－
SC
WC-(Fe, Co, Ni)
RI PT
Specimens
1300 WC-Co
AC C
EP
TE D
1400
18
925±53 －
0.59
－
0.42
683±65
0.28
790±49
M AN U
1350
830±71
Legends
ACCEPTED MANUSCRIPT
Figure 1 SEM photographs of the surface morphology: (a) WC–(Fe–Ni–Co), and (b) WC–Co composite powders, respectively.
Figure 2 XRD patterns of (a) WC–(Fe–Ni–Co), and (b) WC–Co alloys after various sintering
RI PT
temperature processes.
Figure 3 Comparison of the volume shrinkage ratio and apparent porosity of (a) WC–(Fe–Ni–Co),
SC
and (b) WC–Co alloys after various sintering temperature processes.
M AN U
Figure 4 BEI morphology observations of WC–(Fe–Ni–Co) specimens by various sintering temperature treatments: (a) 1250°C, (b) 1300°C, (c) 1350°C, and (d) 1400°C.
Figure 5 BEI morphology observations of WC–Co specimens by various sintering temperature
TE D
treatments: (a) 1250°C, (b) 1300°C, (c) 1350°C, and (d) 1400°C.
Figure 6 Comparison of the mechanical tests of WC–(Fe–Ni–Co) and WC–Co hard metal alloys
EP
after various sintering temperature processes: (a) hardness test, and (b) TRS test.
AC C
Figure 7 Comparison of the fracture toughness of WC–(Fe–Ni–Co) and WC–Co hard metal alloys after various sintering temperature processes.
Figure 8 SEM observations of the (a) 1300°C sintered WC–(Fe–Ni–Co), and (b) 1350°C sintered WC–Co hard metal alloys after KIC tests.
Figure 9 SEM observations of the bridging role of the (a) 1300°C sintered WC–(Fe–Ni–Co), and (b) 1350°C sintered WC–Co hard metal alloys after KIC tests.
19
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
Fig. 1
20
ACCEPTED MANUSCRIPT
Fig. 2
SC
RI PT
(a)
AC C
EP
TE D
M AN U
(b)
The color figure on the Web only 21
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
Fig. 3
The color figure on the Web only 22
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
Fig. 4
23
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
Fig. 5
24
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
Fig. 6
The color figure on the Web only
25
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
Fig. 7
The color figure on the Web only
26
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
Fig. 8
27
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
Fig. 9
28
ACCEPTED MANUSCRIPT
1. The sintered nano WC–Co alloy sintered at 1350°C had the highest hardness (HRA 90.92). 2. The sintered nano WC–(Fe–Ni–Co) alloys showed a good contiguity of 0.44. 3. The optimal nano WC–Co sintered alloy possessed the highest TRS value (2860.08 MPa).
RI PT
4. WC–(Fe–Ni–Co) sintered alloy possessed the highest fracture toughness of KIC (16.23 MPam1/2).
AC C
EP
TE D
M AN U
SC
5. The adding of an iron–nickel instead of a cobalt binder for tungsten carbides is preferable.