Nanometric WC-12 wt% AISI 304 powders obtained by high energy ball milling

Advanced Powder Technology 30 (2019) 1018–1024

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Original Research Paper

Nanometric WC-12 wt% AISI 304 powders obtained by high energy ball milling C.M. Fernandes, J. Puga, A.M.R. Senos ⇑ Department of Materials and Ceramics Engineering, CICECO, University of Aveiro, 3810-193 Aveiro, Portugal

a r t i c l e

i n f o

Article history: Received 6 June 2018 Received in revised form 21 November 2018 Accepted 21 February 2019 Available online 1 March 2019 Keywords: Nanometric powders Tungsten carbide Stainless steel High energy ball milling

a b s t r a c t WC cemented carbides with a greener alternative binder to Co, AISI 304 stainless steel (SS), were processed through high energy ball milling (HEBM). The milling parameters, such as rotation speed, ballto-powder ratio and milling time were investigated. Selected milling conditions were applied to obtain a nanosized powder of WC-12 wt% SS with a highly uniform distribution of the ductile phase. For comparison, a conventionally wet milled powder was also prepared. Both powders were thermally characterized by dilatometry, up to 1450 °C, using vacuum atmosphere, and structural and microstructural analysis were performed in the sintered samples. The nanometric size of the HEBM powder particles markedly affected its densification and thermal reactivity; when compared with the micrometric powder obtained from conventional milling, early starting densification, with a greater contribution of solid state sintering, and increased reactivity, with formation of a larger amount of (M,W)6C phase, was noticed during sintering of HEBM powder compacts. Ó 2019 The Society of Powder Technology Japan. Published by Elsevier B.V. and The Society of Powder Technology Japan. All rights reserved.

1. Introduction Tungsten carbide based composites are considered one of the oldest and most successful powder metallurgy products [1]. The standard tungsten carbide based composites have cobalt as binder, in contents between 4 and 15 wt%. Among others, these materials find applications, as components for mining, oil and gas drilling, transportation and construction, metal forming, structural components and forestry tools [2]. However, due to the cobalt toxicity and high variable market value, efforts are being made targeting its substitution. Stainless steel stands as an interesting alternative to the traditional cobalt binder, since it is more economic, less toxic, and presents competitive technological properties when compared to the standard WC-Co hardmetals [3,4]. The first reported case in the substitution of Co by austenitic stainless steel (AISI 316) in hardmetals was made by Farooq et al. [5]. Since then, efforts have been made aiming to understand and optimize the WC-stainless steel hardmetal system. More recently, results using AISI 304 stainless steel (SS) binder added to WC were reported by Fernandes et al. [6–8] and it was revealed that this binder shows good sinterability characteristics due to the excellent wetting of the molten binder on the WC surfaces [8]. Both

⇑ Corresponding author.

innovative powder coating process and conventional powder metallurgical routes were investigated to process WC-SS cemented carbides [8,9]. One of the challenges in the processing of WC-SS is the composition control, since the stoichiometric carbon content falls inside the M6C + WC + fcc region, as shown in the closest phase diagram of W-C-Fe presented in Fig. 1 [10]. However, the addition of a correct carbon content to WC-SS compositions can avoid the formation of M6C, a mixed carbide phase which is deleterious in terms of mechanical properties [3,11]. The mechanical characterization of these composites with a coarse grain size (
3 mm) shows similar properties to the matching WC-Co compositions. Currently, the need for hardmetals with improved properties, particularly increased hardness and strength, coupled with high toughness, have challenging researchers to develop increasingly finer grained hardmetals [12,13]. Despite of the numerous techniques to produce nanostructured materials, high-energy ball milling (HEBM), also known as mechanical alloying, has become the most popular method to fabricate nanocrystalline materials due to its simplicity, relatively inexpensive equipment and its potential for large-scale production [14]. HEBM is fundamentally different from the conventional ball milling (CM), the traditional powder processing technique used in the hardmetal industry. The major differences between those milling processes are found on the impact energy (HEBM is typically 1000 times higher than

E-mail address: [email protected] (A.M.R. Senos). https://doi.org/10.1016/j.apt.2019.02.016 0921-8831/Ó 2019 The Society of Powder Technology Japan. Published by Elsevier B.V. and The Society of Powder Technology Japan. All rights reserved.

C.M. Fernandes et al. / Advanced Powder Technology 30 (2019) 1018–1024

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For comparison, the WC-12SS composite powder was also blended in a planetary ball mill, hereafter refereed as conventional milling (CM), with WC-Co balls for 8 and 20 h, at 200 rpm, using 2propanol as milling media (Table 1). The median particle size, G50, was measured by dynamic light scattering (DLS) analysis with a Zetasizer Nano Zs – Malvern Instruments. The water based suspensions were prepared with 0.02 wt% of powders ultrasonicated for 15 min. The structural characterization was performed in a Rigaku diffractometer with Cu Ka radiation, while the crystallite size was achieved through the Scherrer equation [16]:

DScherrer ¼

Fig. 1. Vertical section of the Fe-W-C phase diagram calculated for 10 wt% of Fe [10]. The solid symbol on the composition axis indicates the stoichiometric composition.

CM), milling action (CM is dominated by attrition, while impact is the main milling action in HEBM), milling time (longer for HEBM), milling atmosphere (controlled in the case of HEBM, using vacuum or gases) and achieved particle size (normally micrometric sizes in CM against nanometric in HEBM) [15]. Besides that, phase’s transformations and chemical reactions can be activated by HEBM [15]. Taking this into account, this work investigates the high energy ball milling conditions to attain nanometric WC-12 wt% SS powders; additionally, the thermal shrinkage behavior of the nanosized powder and the structural and microstructural characteristics of the respective densified composites were assessed and compared with the same features in composites obtained by conventional milling.

2. Experimental procedure Tungsten carbide (WC) powder (H.C. Starck, HCST-Germany) and AISI 304 stainless steel (SS) powder (Sandvik Ospery) with an average particle size of 1 lm and 5 lm, respectively, were used. The approximate chemical composition of the SS powder, given by the supplier, is 68.2 Fe; 18.8 Cr; 10.0 Ni; 2.0 Mn; 1.0 Si and 0.03 C (wt%.). The starting WC powder is constituted by irregular shaped particle aggregates (Fig. 2a), while the AISI 304 SS powder has coarser particles with near spherical morphology (Fig. 2b). The WC powder presents only the WC phase, as depicted in Fig. 2c, and the AISI 304 SS powder presents two iron rich phases, the main is austenitic (Fe-c) and the minor martensitic (Fe-a), Fig. 2d. The high energy ball milling (HEBM) of WC with 12 wt% of SS metallic binder (WC-12SS) was performed in a planetary mill Fritsch Pulverisette 6, using a stainless steel (AISI 303) bowl and WC-Co balls, Table 1. The argon atmosphere is used to avoid the powder and grinding media oxidation during milling. To attain a proper argon atmosphere, the grinding bowl was purged during 2 min and pressurized at 200 kPa (2 bar). The investigated milling parameters were the rotation speed, ball-to-powder ratio (BPR) and milling time, according to Table 1. HEBM of individual WC powders were also studied, under some milling conditions, to understand the role of the ductile phase on the milling process.

Kk b cos h

ð1Þ

where DScherrer is the weighted apparent crystallite size, b the integral breath of the line profile (XRD peak), 2h the diffraction angle, k the X-ray wavelength (0.154056 nm for Cu Ka1 radiation) and K the Scherrer constant (shape-dependent, being 0.89 for spherical crystallites) [16]. The experimental error in the determined crystallite values was 10%. The powders were uniaxially pressed (CARVER laboratory press – Model C) with a cylindrical shape (10 mm diameter and 3 mm thickness) during 60 s up to 250 MPa, followed by cold isostatic pressing (High pressure system U33, Institute of High Pressure Physics) at 350 MPa, during 15 min. The sintering cycle was performed in a vacuum furnace at 1450 °C, during 1 h at 30 Pa, with a heating/cooling rate of 10 °C/min. The microstructural characterization was assessed by scanning electron microscopy (SEM, Hitachi SU-70) with energy dispersive X-ray spectroscopy (EDS, Bruker-Quantax 400). The carbon content of composite powders was determined, after milling and after the thermal consolidation, by automatic direct combustion equipment (LECO CS 200 IH). The theoretical density of the WC-12SS composition was calculated by the mixing rule, assuming dSS = 7.93 g/cm3 and dWC = 15.58 g/ cm3. The thermal reactivity of the composite powders was evaluated by differential thermal analysis (DTA) using a SETARAM Labsys equipment and by dilatometry using a homemade vertical graphite dilatometer. Parallelepipedic samples (14  4  4 mm3) of HEBM and CM powders were uniaxially and isostatically pressed at 200 MPa and 350 MPa, respectively. The shrinkage of the pressed compacts was measured from room temperature up to 1450 °C, with a constant heating rate (CHR) of 10 °C/min under vacuum atmosphere (1 Pa). The measured shrinkage was corrected for the dimensional variation of the equipment upon heating, using the length variation of a graphite sample submitted to the same heating schedule. DTA tests of HEBM and CM powders were performed, as well, at CHR of 10 °C/min up to 1400 °C, in argon atmosphere. 3. Results and discussion 3.1. HEBM optimization In Fig. 3a, the XRD patterns of WC-SS powder prepared by HEBM, at 350 rpm and BPR 20:1, illustrate the effect of the time variation, between 2 and 20 h, in the powder structure. It is noticeable a peak broadening with increasing milling time which is more relevant in the first 2 h of milling. This peak broadening results from the grain refinement and/or micro-strain generated during the commination step [17]. In consequence, the peaks correspondent to the secondary iron rich phases can’t be detected in the HEBM milled powders, only the ones of the WC major phase. This is not the case of WC-12SS powder conventionally milled, CM, as shown in the XRD of Fig. 3b: the peaks of Fe(c) and Fe(a) phases, already detected in SS original powder, Fig. 2d, are here clearly discernible. As referred before, the HEBM milling process induces the

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a)

b)

30 µm

m 3 µm

d)

WC

WC

30

Fe (γ)

WC Intensity (a.u.)

Intensityy ((a.u.))

c)

34

38

42

46

50

54

38

Fe (γ)

Fe (α)

40

42

44

2θ (º)

46 4 8 2θ (º)

50

52

54

Fig. 2. SEM micrographs of the raw powders: (a) WC and (b) AISI 304 SS; and respective XRD diffraction patterns: (c) WC and (d) AISI 304 SS.

Table 1 Experimental conditions for high energy ball milling, HEBM, and conventional milling, CM. CM

Bowl composition Balls composition Milling media Rotation speed (rpm) Ball-to-powder ratio (BPR) Milling time (h)

AISI303 WC-Co dry milling 200, 350 10:1, 20:1 2, 4, 6, 8, 10, 20

AISI303 WC-Co 2-propanol 200 4:1 8, 20

crystal structure refinement and, consequently, the broadening of the XRD peaks, reducing the detection limit of minor phases, as is the case of the iron rich phases. Therefore, the calculation of the crystallite size of HEBM milled powders from the XRD patterns, using Eq. (1), could only be done for the WC main phase. Fig. 4 presents the variation of the crystallite size of WC-12SS powder with the HEBM time at different conditions of rotation speed and BPR. The results show the expected decrease of the crystallite size with the milling time for constant rotation speed and BPR. Besides, it can also be observed that the increase of either the milling rotation speed, from 200 to 350 rpm, or the BPR, from 10:1 to 20:1, induces the decrease of the WC crystallite size for the same milling time. Furthermore, the powders submitted to a rotation speed of 200 rpm only attained a minimum crystallite size of 14 nm, for 20 h of milling time and BPR = 20:1, whereas for the same BPR and using 350 rpm, a reduced value of
12 nm can be attained for only 2 h of milling, and this crystallite size will be stabilized in
11 nm after 4 h of milling time. Therefore, taking the effect of the rotation speed and BPR conditions on the crystallite size, we can establish that the most efficient HEBM conditions for WC-12SS composites were observed for 350 rpm and BPR = 20:1.

20h

Intensity (a.u.)

HEBM

30

10h 8h 6h 4h 2h WC

32

34

36

38

40

42 44 2 (°)

46

48

50

52

54

(a) WC WC

Intensity (a.u.)

Conditions

WC

HEBM-8h Fe(γ) CM-8h

Fe(γ) Fe(α)

30

32

34

36

38

40

42 44 2 (°)

46

48

50

52

54

(b) Fig. 3. (a) Diffraction patterns of the WC-12SS powders with different milling times, at 350 rpm with a ball-to-powder weight ratio of 20:1, showing peak broadening with increasing milling times. The diffraction pattern of the precursor WC powder is added for comparison. (b) XRD patterns of WC-12SS composition, after 8 h of milling in HEBM (350 rpm and BPR 20:1) and in CM (200 rpm, BPR 4:1).

Particle size measurements were performed with DLS technique in the WC and WC-12SS powders milled under the optimized conditions of 350 rpm and BPR 20:1 and using variable milling time

C.M. Fernandes et al. / Advanced Powder Technology 30 (2019) 1018–1024

40 200 RPM - 10:1 200 RPM - 20:1 350 RPM - 10:1 350 RPM - 20:1

Crystalite size (nm)

35 30 25 20 15 10 5 0 0

2

4

6

8

10

12

14

16

18

20

22

24

Time (h) Fig. 4. WC crystallite size under different milling conditions for the WC-12SS composite powder. The dashed circle on the crystallite size axis indicates the initial value for the WC powder.

1021

neous distribution of the Fe element is well achieved in the HEBM process, as depicted in Fig. 6a. Comparing the Fe distribution on the WC-12SS powders milled by HEBM in the optimized conditions (350 rpm, BPR 20:1, 8 h) and by CM (200 rpm, 20 h) (Fig. 6b), it is well perceptible the much higher homogenization achieved in HEBM powders, together with a finer particle size. These characteristics, i.e., the nanometric particle size and the high homogeneity of the ductile phase distribution attained in the HEBM composite powders, are very important factors from the technological point of view, since they have a great impact on the powder processing, final microstructure and, therefore, on the final composite properties. Taking this in mind, the thermal behavior and the microstructure of HEBM sintered samples will be further investigated in this work and compared with the correspondent characteristics of CM powder samples. 3.2. Thermal behavior

(Fig. 5). An effective decrease of the average particle size, from
1100 nm to c.a. 100 nm, can be observed for the WC powder after 2 h of milling time, but no appreciable variation is detected for longer milling times (Fig. 5). On the other hand, the behavior of the composite powder, constituted by brittle (WC) and ductile SS particles, is slightly different from that of the single phase WC powder. In the initial milling stages, up to 4 h, there is a decrement of the particle size to
144 nm in consequence of WC particles fragmentation and metallic particles flattening by the ballpowder-ball collisions, as was firstly reported by Benjamin et al. [18]. With further milling, the ductile particles get work hardened, acquiring a lamellar shape and agglutinating the brittle particles between them, which results in coarser particle size distributions [19]. This coarsening effect is observed to occur during HEBM of the WC-12SS powder, in Fig. 5, where an increase of the average particle size from
140 nm up to
230 nm, occurs between 4 h and 8 h of milling time. For longer milling times the particle size keeps constant. The milling time of the WC-12SS powder was, then, selected taking into account both the results of crystallite size and particle size. As pointed before, after 4 h of milling, at 350 rpm and 20:1 BPR, there was no appreciable change in the WC crystallite size. On the other hand, under the same conditions of rotation speed and BPR, the average particle size showed, after the initial decrement, an increase between 4 and 8 h, caused by the covering of the brittle fragments with the ductile SS phase, which may correspond to a higher uniformity of the SS binder distribution. Therefore, 8 h of milling time was the selected parameter to get a nanometric WC powder uniformly coated by the ductile SS phase. In order to analyze the morphology of the SS binder and the chemical distribution of the metallic elements, Fe-EDX maps were performed in HEBM and CM composite powders. The homoge-

Particle size (nm)

1200 1000 800 600 400 WC-12SS

200

WC

0 0

2

4

6

8

10 12 14 Time (h)

16

18

20

22

24

Fig. 5. Particle size evolution during HEBM of WC and WC-12SS powders (350 rpm, 20:1 BPR). The dashed circle on the particle size axis indicates the initial value for the WC powder.

The thermal behavior of the WC-12SS composite powders in terms of reactive sintering was studied by dilatometry and DTA analysis, performed with the optimized HEBM powder (350 rpm, BPR 20:1, 8 h) and the CM powder (200 rpm, 20 h), for comparison. The sintering behavior of WC-SS powders with varied SS binder contents was recently reported [8] and three densification stages were identified, as can also be observed in Fig. 7, for the CM powder: (I) initial slower shrinkage in solid state sintering, starting
900 °C up to
1180 °C; (II) a second stage with a much higher densification rate governed by viscous flow mechanisms and (III) a third and final stage with a decreased rate of densification and controlled by diffusion mechanisms assisted by a liquid phase. On the other hand, the linear shrinkage and shrinkage rate versus temperature curves of the HEBM powder, in Fig. 7a and b, are quite different: the shrinkage starts earlier,
600 °C, and only two densification stages with different characteristics were perceptible: (i) initial stage, extending from a large range of temperatures,
600 °C up to
1320 °C, with a low shrinkage rate, but attaining at the end a significant 6% of shrinkage; (ii) a second stage, for higher temperatures, with a fast acceleration of the shrinkage. The observed total shrinkage, Y = 14%, at 1450 °C (limiting temperature of the equipment) corresponds to
87% of relative density (Table 2). However, it can be observed that the shrinkage rate at that temperature is yet high enough to preview the good progress of the densification using higher temperatures or a holding time at 1450 °C. In fact, 99% of relative density could be attained with a holding time of 30 min at 1450 °C (Table 2) [20]. Looking at the DTA analysis of HEBM and CM powders, in Fig. 8, it can also be seen some differences in the respective heat flow vs temperature curves. Whereas for the CM powder two endothermic events are detected at
860 °C and
1188 °C, which were identified with the formation of a mixed carbide phase (eta-phase) and an iron rich liquid phase, correspondingly [8,11], only one endothermic peak can be detected at 595 °C for the HEBM powder. The XRD diffraction patterns of the initial HEBM powder and after a thermal treatment at 600 °C, during 60 min, in Fig. 9, show that eta-phase is clearly detected after the thermal step. The endothermic peak of 595 °C may, then, correspond to the early formation of the mixed carbide phase in the nanometric HEBM powder at lower temperature than have been observed in micrometric WC-SS powders, T > 800 °C [8,11], as well as in the CM powder (G50
1 lm, Table 2), T
860 °C. Besides, a noticeable content of eta-phase was already formed in the thermal treated HEBM powder,
22 wt%, as calculated by the Rietveld analysis of the respective XRD pattern, presented in Fig. 9. The eta-phase formation is thermodynamically favored using iron rich binders as SS [11,21] and its content increases with decreasing the carbon content to values lower than the stoichio-

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Fig. 6. SEM micrographs and respective Fe X-ray map of WC-12SS milled powders at (a) HEBM (350 rpm, BPR 20:1, 8 h); (b) CM (200 rpm, 20 h).

18 (a)

16

70

CM

III

60 Heat Flow (mW)

HEBM

14 12 Y(%)

10 II

8

II

6 4

I

CM HEBM

50 40

Endo

1188ºC

30 20 10 857oC

0

2

I

-10

0 500

700

900 1100 Temp. (ºC)

1300

1500

595ºC

-20 500

700

16

dY/dt (10-4 s-1)

CM

(b)

14

900 1100 Temperature (ºC)

1300

1500

Fig. 8. DTA analysis performed in argon atmosphere, at 10 °C min1, for HEBM and CM powders.

12 10

during HEBM, lowering the carbon content to c.a. 4.98 wt%. The decarburization is yet more pronounced during the heating up to 1450 °C, and a reduced C amount of c.a. 3.81 wt% was determined in the sintered sample (Table 2). On the other hand, the decarburization in the CM powder is negligible and, after heating up to 1450 °C, a higher amount of C,
5.08 wt% remains in the sintered sample. Taking these results into account, the shrinkage stages of the HEBM dilatometric curve can be further interpreted: (i) Stage I, starting at temperatures close to the eta-phase formation, is controlled by diffusion mechanisms of solid state sintering, as in the initial stage of CM powders; (ii) Stage II, for temperatures higher than 1320 °C, and higher shrinkage rates will correspond to the actuation of liquid phase sintering mechanisms. In fact, the strong decarburization observed in HEBM powders during milling and heating will deviate the composition for the left side in the FeW-C phase diagram, Fig. 1, to the WC + M6C field, and the formation of a reduced amount of liquid phase is only expected for tem-

8 6 HEBM 4 2 0 500

700

900 1100 Temp. (ºC)

1300

1500

Fig. 7. (a) Linear shrinkage (Y) and (b) shrinkage rate (dY/dt) vs temperature, heated at 10 °C min1 for HEBM and CM powders.

metric ones. The carbon content of the milled powders and samples sintered up to 1450 °C was determined and is presented in Table 2. Taking the value of 5.39 wt% for the stoichiometric carbon in the WC-12SS composition (correspondent to 6.13 wt% for WC), it can be observed that there is already a significant decarburization

Table 2 Physical, chemical and microstructural characteristics of WC-12SS samples before and after sintering. Sample

CM HEBM

Green

Sintered

G50 (nm)

C0 (wt%)

d0 (%)

Y1450 (%)

d1450 (%)

d1450,

980 230

5.30 4.98

55 55

15.9 14.1

93 87

99 99

30 min

(%)

Cs (wt%)

M6 C M 6 CþWC

5.08 3.81

11 48

(wt%)

G (nm) 1200 151

G50 - median particle size; C0 - carbon content after milling; d0 - relative green density; Y - Shrinkage; d - relative density; Cs - carbon content after sintering; G - average grain size.

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Intensity (a.u.)

WC

WC

WC

η

η

η 600ºC, 1h After milling

30 32 34 36 38 40 42 44 46 48 50 52 54 56 58 60 2 teta (º) Fig. 9. XRD patterns of WC-12SS powders after HEBM and after thermal treatment at 600 °C.

WC

WC

Intensity (a.u.)

WC

η η

η

η

η

γ

η

HEBM

α CM 30

35

40 2 teta (º)

45

50

Fig. 10. XRD diffraction profiles of samples sintered at 1450 °C, from CM and HEBM powders. (g-phase represents Fe3W3C phase).

(a)

peratures higher than 1280 °C. The reduced content of formed liquid can also explain the absence of an endothermic peak in the respective DTA curve, Fig. 8. The structure and microstructure of the sintered samples were analyzed by XRD and SEM analysis. The XRD diffraction profiles of the samples after sintering at 1450 °C (Fig. 10), indicate the presence of the major WC phase, eta-phase (M6C phase, where M represents essentially Fe, W, Cr) and a minor iron rich phase. The calculated percentage of M6C is 11 wt% for CM sample and a huge amount of 48 wt% could be detected in the HEBM sample, caused by the stronger decarburization of the highly reactive nanometric powders. A similar increased reactivity was already observed for the WC powders sputter-coated with nanometric SS binder [11]. In consequence, almost all the SS binder elements react with WC to form eta-phase in the HEBM composition and only <2 wt% is estimated to be available to form liquid phase at the sintering temperatures. In accordance to this, the XRD of the HEBM sample presents only a vestigial peak of an austenitic iron rich phase, Fe(c), whereas a larger peak of ferritic/marthensitic phase, Fe(a0 ), could be detected in the CM spectrum, Fig. 10. Interesting to note that the higher formation of eta-phase stabilizes the austenitic phase, probably by leaving a higher Ni content available to form the iron rich phase (Ni can be accommodated in the M6C structure but in reduced percentage, since is not a carbide former as Fe and Cr [11]). Regarding the SEM micrographs of CM and HEBM composites after sintering (Fig. 11), two major phases can be easily identified; WC grains in light gray colour and eta-phase in dark gray colour. For the HEBM sample, nanometric WC grains (
150 nm, Table 2) embedded in large regions of eta-phase are noticeable in the SEM microstructure (Fig. 11a and b), while smaller regions of eta-phase well dispersed between micrometric WC grains (
1.2 lm of mean grain size, Table 2) can be observed for the CM sample (Fig. 11c). No abnormal grain growth was observed to occur in both cases. Moreover, a dark phase attributed to the SS

(b)

Eta-phase

5µm

2.5µm

(c)

5 µm Fig. 11. SEM micrographs of sintered samples (10 °C min1 until 1450 °C): (a) and (b) HEBM and (c) CM.

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binder (Fe(a) peak in Fig. 11) is also observed among the WC grains of the CM composite. The nanometric structure of the composites produced by HEBM envisage interesting mechanical properties, regarding hardness [19,20]; on the other hand, an adequate carbon addition to the precursor powders to compensate the decarburization along the process can be used to increase other properties, such as the mechanical strength and ductility [22]. 4. Conclusions Nanostructured powders of WC cemented carbide with a fixed binder content of 12 wt% AISI 304 stainless steel, WC-12SS, were produced by high energy ball milling (HEBM). An efficient reduction of crystallite and particle size was observed for all the tested conditions of rotation speed, ball to powder ratio (BPR) and milling time. The increase of rotation speed from 200 to 350 rpm had a greater influence on the reduction of the crystallite size than the increase of BPR from 10:1 to 20:1. A minimum crystallite size of
11 nm could be achieved with 350 rpm and 20:1 BPR, after 4 h of milling time, but increasing the milling time up to 8 h improved the binder distribution due to the spreading of the ductile SS phase on WC particles. By applying the optimized HEBM parameters, a nanometric powder with a quite homogeneous distribution of the iron rich phase, as confirmed by Fe-EDX maps, could be obtained. The densification and thermal reactivity of the HEBM powder presented a distinct behavior from that observed for a conventionally mixed (CM) micrometric powder with the same nominal composition. The shrinkage of HEBM powder started earlier,
600 °C, and only two densification stages could be observed, instead of three stages, after
900 °C, in the CM powder. This was interpreted as a favored solid state sintering in the HEBM powder, coming from its nanometric scale and large decarburization during milling and heating. These characteristics also triggered the formation during heating of a larger amount of (M,W)6C phase. High dense composites, keeping the nanometric structure of the WC grains, could be obtained at sintering temperatures
1450 °C. Acknowledgements The authors wish to thank Prof. Dr. L.F. Malheiros and MSc C. Lopes for the assistance in DTA/TG analysis. The work was developed within the scope of the project CICECO-Aveiro Institute of Materials, POCI-01-0145-FEDER-007679 (FCT Ref. UID/ CTM/50011/2013), financed by national funds through the FCT/ MEC and when appropriate co-financed by FEDER under the PT2020 Partnership Agreement.

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