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Home All issues Volume 111 / No 3 (2023) Matériaux & Techniques, 111 3 (2023) 304 Full HTML
Open Access
Issue
Matériaux & Techniques
Volume 111, Number 3, 2023
Article Number 304
Number of page(s) 13
Section Materials production and processing
DOI https://doi.org/10.1051/mattech/2023016
Published online 28 June 2023

© K. Papy et al., 2023

Licence Creative CommonsThis is an Open Access article distributed under the terms of the Creative Commons Attribution License (https://creativecommons.org/licenses/by/4.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

1 Introduction

Laser power bed fusion (L-PBF), also known as selective laser melting (SLM), is an additive manufacturing (AM) process used for the production of functional parts with complex shape using a metallic powder as feedstock material. In L-PBF a laser beam melts a powder layer, which after solidification builds up the three-dimensional part layer-by-layer [1–3]. The final physical properties such as roughness, density, porosity, cracking propensity depend on process parameters including the laser power, scanning speed, and powder bed thickness [4–7].

Tungsten carbide/cobalt (WC-Co) cermets are extensively used in the manufacture of cutting tools. The properties and applicability of tungsten alloys depend significantly on the WC content (from 80 to 97 wt.%) as well as the binder phase composition. In general WC-Co cermets exhibit exceptional hardness, low friction coefficients, and high melting points, all of which contribute to a high wear resistance.

WC-Co parts are classically manufactured by powder sintering [8–12]. The layer-by-layer AM method could significantly increase tool design complexity and potentially improve tool performance. Several studies regarding L-PBF manufacturing of WC-Co parts are available in the literature [13–17]. Generally, it is reported that manufacturing good quality parts is challenging due to the high melting point, high thermal conductivity, and brittle nature of the as-built parts. In most cases their integrity was low owing to the presence of long cracks and pores. For example, Uhlmann et al. [18] reported that WC-17Co parts manufactured by L-PBF at low (185 J/mm3) and high (1667 J/mm3) energy densities exhibited substantial porosity and cracking. They also established that parameters such as beam focalization, layer thickness, laser power, and scan speed affected the relative density of the WC-17Co material. In contrast to conventional sintering, L-PBF process used laser energy, which lead to a thermal exposure with rapid solidification. This makes the process prone to the formation of pores, cracks and brittle phases with cermet material such as WC-Co. The work provides an interesting optimization study of process parameters, but the possible benefit of post heat-treatment is not explored. Khmyrov et al. [19] used a mixture of WC-75Co and WC-50Co powders to determine the optimal WC content at which cracking can be avoided, mainly caused by the presence of the ternary Co3W3C brittle phase. It was found that cracking of AM parts could be avoided only at low percentages of WC, which is not interesting for applications (compared to cermets obtained by sintering). The main objective of the study conducted by Schwanekamp et al. [20–22] was to reduce the porosity and crack frequency by preheating the support plate. Preheating to 800 °C resulted in a significant reduction in crack prevalence and porosity, but a large number of micropores still remained. Further post-treatment by hot isostatic pressing (HIP) resulted in a significant improvement of the material density. The work concluded that the use of a preheated building plate was essential for the successful manufacture of WC-Co parts using L-PBF. The positive effect of HIP reported in their work is encouraging. Therefore, further studies exploring its applicability limits in terms of the microstructure of AM parts could help developing energy efficient methodologies.

Considering the above literature results it can be concluded that L-PBF manufacturing of WC-Co parts remains challenging. Preheating of the building plate to a high temperature allowed the manufacturing of solid samples with some residual porosity that could be eliminated through HIP post-treatment. However, in the absence of a preheating plate the manufacturing of complex-shaped WC-Co parts, even with high internal porosity, was not possible owing to crack formation and to the destruction of the part during cutting from the plate. In addition, it should be confirmed whether the HIP process is able or not to heal the defects formed during AM. Additional information concerning microstructure changes during HIP would be also helpful to understand final physical properties of the manufactured parts.

Present article focuses on the AM of parts using the commercial WC-17Co material and addresses three essential points as illustrated in Figure 1. First, the possibility to produce complex shapes by L-PBF without using a preheated plate will be tested and the optimal process parameters will be determined. In a second step the actual benefits of the HIP post process will be evaluated. Especially a comparison with a conventional heat treatment will demonstrate the importance of the applied pressure. Finally, the hardness performance of the produced parts will be tested.

thumbnail Fig. 1

Global scheme of the study with three main steps: L-PBF process, post heat-treatment and hardness tests.

Schéma descriptif de l’étude en trois étapes principales : procédé L-PBF, post-traitement thermique et essais de dureté.

2 Material and methods

2.1 L-PBF Equipment

Pro X 200 DMP equipment supplied by 3D Systems was used in the study. The volume of the working chamber where the parts were manufactured was 140 × 140 × 125 mm3. The L-BF machine was equipped with a 400 W fiber laser having a wavelength of 1070 nm. To prevent oxidation the working chamber was filled with nitrogen during manufacturing.

It is important to note that the analysis of the laser beam shape revealed a significant effect on the energy distribution in the beam. In particular, at powers higher than 200 W and long exposure times the energy distribution changed from a Gaussian spot to a ring shape. The energy distribution surfaces measured at laser powers of 130 W and 300 W are presented in Figure 2. Evidently, this shape change significantly complicates the optimization process. To avoid its influence the maximum laser power was limited to 160 W so that the shape of the energy distribution surface remained close to Gaussian. The spot diameter was 130 μm.

thumbnail Fig. 2

Intensity distribution surfaces at (a) 130 W and at (b) 300 W (laser power in the focal plane).

Distribution de l’intensité laser à (a) 130 W et à (b) 300 W (au plan focal).

2.2 Powder material

Agglomerated-sintered WC-17Co powder with a particle size distribution of 15–53 μm (Oerlikon Metco) was used in this study. The powder was initially developed for thermal spray applications such as HVOF and HVAF. SEM images of the particles are shown in Figure 3. It is evident that the particles are quasi-spherical in shape.

The average size of the WC grains (dark grey) is in the 1–5 μm range. The carbide grains are uniformly distributed in the cobalt matrix (light grey).

Inductively coupled plasma mass spectrometry (ICP-MS) analysis (see Tab. 1) indicates that the chemical composition was within the range stated by the supplier (C: 4.7–5.5, Co: 14.5–19.5).

X-ray diffraction (XRD) was employed for phase analysis using an X’PertPro MPD diffractometer with CuKα radiation. Phase quantities were obtained from Rietveld analysis performed with the MAUD software [23]. The powder pattern shown in Figure 4 was fitted with an Rwp = 4.1% (χ2 = 3.4), which resulted in mass fractions of 90 ± 6% for the WC phase (space group P-6 m2) and 10 ± 4% for the Co phase (space group Fm-3m). The corresponding calculated volume fraction is 0.84% and 0.16% for the WC and Co phases respectively. The phase fraction errors are relatively large, and the results correspond to the lower limit of the nominal composition stated by the manufacturer. The crystallite sizes in the WC and Co phases were 340 and 210 nm, respectively. Additionally, the microstrain in the Co phase was large (∼ 2 × 10−3), whereas it was minimal in the WC phase (∼ 4 × 10−4). The difference between the ICP-MS and XRD results may be related to differences in micro absorption and crystallite size of the two phases.

Various particle properties, such as shape, size distribution, flowability, apparent density and tap density were also analyzed. Powder morphology analysis indicated that the particles were largely spherical with an equivalent diameter of d10 = 24.3 μm, d50 = 47.8 μm and d90 = 73.4 μm. The powder flowability results were obtained using a Hall flowmeter in accordance with ASTM B527 and ASTM B21, and are presented in Table 2 together with the tap and apparent densities.

thumbnail Fig. 3

SEM images of the WC-17Co particles at different magnification (a) x400 and (b) x3000.

Images MEB de la poudre de WC-17Co suivant différents grossissement (a) x400 et (b) x3000.
Table 1

Chemical analysis of the initial powder according to ICP-MS.

Analyse chimique de la poudre initiale par procédé ICP-MS.

thumbnail Fig. 4

XRD pattern of the WC-17Co powder used for L-PBF process.

Diffractogramme de la poudre de WC-17Co utilisée pour le processus L-PBF.
Table 2

Summary of powder properties.

Résumé des propriétés de la poudre.

2.3 Post-treatment processes

Cube-shaped samples with dimensions of 10 × 10 × 10 mm3 manufactured using L-PBF were used for post-treatment. The first batch of samples was heat-treated (HT) in an oven at 1450 °C under nitrogen atmosphere for 5 h. The second batch was subjected to capsule-free HIP, which is a cycle applied during industrial processes. The HIP process was performed in several steps. First, the temperature was gradually increased from room temperature to 1000 °C over 15 h under normal pressure, then it was then elevated to 1450 °C at a heating rate of 5 °C/min. The furnace was maintained at a pressure of 40 MPa for 5 h. Cooling to room temperature was performed in the HIP chamber at a rate of 6 °C/min.

2.4 Structural and mechanical properties

The relative density of the manufactured parts was determined employing the Archimedes method using ASTM-B962 as well as the analysis of scanning electron microscope (SEM) images (using ImageJ [24]). The microstructure and elemental distribution were analyzed using a Zeiss Supra 55VP v2 SEM. The mean grain intercept method is a technique used to quantify the grain size of a given material by drawing a set of randomly positioned line segments on the micrograph and then measuring the length of each segment that intersects a grain [24–26]. Vickers hardness (HV) was determined by applying a load of 30 kg for 10 s to all samples. Palmqvist toughness was determined by cracks propagation measurement after indentation test [27,28].

3 Results and discussion

3.1 Optimization of process parameters

In the first step, the parametric window of the L-PBF process was determined. The principal purpose of these experiments was to determine the parameters enabling the production of materials with the lowest amount of porosity and cracks. Initially, single tracks were deposited with a powder layer thickness of 30 μm. The power and scan velocity were varied from 100 to 160 W and from 30 to 200 mm/s, respectively. Table 3 summarizes the parameter ranges tested in the first stage.

Observation of the track shape allowed to conclude that stable tracks were formed only when the scanning speed was in the 30–40 mm/s range at a laser power of 130 W. Following the single-track experiments, 3D objects with simple geometries (cubes and parallelepipeds) were manufactured. Numerous factors can affect the process results, such as the laser power, scan speed, hatch spacing, layer thickness, and scanning strategy. These factors are important because re-melting depends on the amount of the energy density (Ev) absorbed during the process [29], which is defined by the following equation:

(1)

where Ev is the energy density (J/mm3), P is the laser power (W), Vs is the scanning speed (mm/s), Hd is the hatch distance (mm) and l is the layer thickness (mm). Similar approaches based on energy density optimization in the L-PBF process have already been successfully applied several times [30,31]. Furthermore, hexagonal and zigzag scanning strategies were applied to produce WC-Co cubes (Fig. 5).

To analyze porosity and other defects the samples were cut parallel to the building direction. Figure 6 (bottom) shows the relationship between sample density and energy density. The plot in Figure 6 (top) indicates that samples with the higher density (98.9%) were manufactured using an energy input of 602 J/mm3 and the hexagonal scanning strategy (Fig. 6b). The samples manufactured at a lower energy density (hexagonal scanning strategy) contained significantly more pores and with larger sizes (Fig. 6a). This dramatic increase in porosity is related to the lack of fusion of the WC-Co powder. An increase in the energy density above 602 J/mm3 led to crack formation in addition of lacks of fusion. The samples manufactured at 833 J/mm3 using the hexagonal scanning strategy contained more cracks than low Ev and had the lowest density, which indicates a high level of residual stress. The optimum energy density for manufacturing in the zigzag scanning strategy was found to be 560 J/mm3. Therefore, due to manufacturing difficulties to produced large part, hexagonal strategy with an optimal Ev of 602 J/mm3 was used. For comparison, the optimal Ev level of the Co–28Cr–6Mo alloy is appreciably lower (∼ 200 J/mm3) and its porosity is close to zero [32].

The density of the zig-zag-generated samples was lower than for the hexagonal scanning strategy, therefore, the final batch was produced using the latter. An image of the building plate and the manufactured samples is shown in Figure 7. Blocks (50 × 25 × 10 mm3) were built for abrasive wear tests, and cubes (10 × 10 × 10 mm3) and cylinders (Ø20 × 10 mm) for various characterizations. Lattice-structured cubes were also built to test the limits of manufacturing using cermet powder.

Table 3

Ranges of L-PBF manufacturing parameters applied in the experiments.

Plages des paramètres du procédé L-PBF appliqués pour les optimisations paramétriques.

thumbnail Fig. 5

Schematic illustration of (a) zigzag and (b) hexagonal scanning strategies.

Illustration schématique des stratégies de balayage (a) en aller/retour et (b) hexagonale.
thumbnail Fig. 6

Top: SEM images of samples manufactured via the hexagonal scanning strategy at (a) low, (b) medium, and (c) high energy densities. Bottom: Plot of sample density vs. energy density for zigzag and hexagonal scanning strategies.

En haut : images microscopes d’échantillons fabriqués selon la stratégie de balayage hexagonale à (a) faible, (b) moyenne et (c) haute densité d’énergie. En bas : évolution de la densité de l’échantillon en fonction de la densité d’énergie volumique pour les stratégies de balayage en aller/retour et hexagonale.
thumbnail Fig. 7

Image of the building plate and the samples manufactured for various analyses.

Image d’un plateau d’échantillons fabriqués pour diverses analyses.

3.2 Microstructure and composition of the as-built samples

SEM images of the as-built sample produced at 602 J/mm3, applying the hexagonal manufacturing strategy, are shown in Figure 8.

Microstructure observations revealed that unlike the powder, the as-built samples contained significantly irregular carbide grains size. Due to laser impact, large grain bursts into smaller grains in the material and smaller grain are diffused in the binder [33]. A mean WC grain size of 5.5 μm was measured. Some submicron carbide inclusions were also present in the metal matrix. Short cracks and pores were visible throughout the sample (Fig. 9), but they did not affect the integrity of the cube.

The pore size distribution evaluated with ImageJ is shown in Figure 10. Ten different micrographs were recorded at randomly selected positions on each cross section using a magnification of x50. Pores with sizes of < 5 μm were uniformly distributed in the sample. Larger pores were present mostly at the interfaces between the connected tracks and layers, indicating that the material was only partially melted in these areas.

The XRD pattern corresponding to the as-built sample is shown in Figure 11. Compared to Figure 5 (containing the WC and Co phases only) complex carbides formed during the L-PBF process are also visible, caused by decarburization phenomena. In particular, the W2C phase and the ternary Co3W3C and Co3W9C4 phases were present. These phases have compact hexagonal (P63/mmc), cubic (Fd-3m), and compact hexagonal structures (P63/mmc), respectively.

Formation of the fragile W2C and Co3W3C phases promote sample cracking during the L-PBF process [34,35]. In conventional powder sintering the W2C phase is considered undesirable and should be avoided.

thumbnail Fig. 8

SEM images of the as-built sample at (a) x500, (b) x1000, and (c) x5000 magnifications.

Images MEB de l’échantillon brut de fabrication aux grossissements (a) x500, (b) x1000, et (c) x5000.
thumbnail Fig. 9

Secondary electron SEM image of the as-built sample illustrating the presence of cracks, pores, and fine and coarse carbide grains.

Image MEB de l’échantillon brut de fabrication, illustrant la présence de fissures, de pores et de différentes tailles de grains de carbure.
thumbnail Fig. 10

Pore size distribution in the as-built samples and after heat treatment.

Distribution de la taille des pores dans les échantillons brut de fabrication et les échantillons après traitements thermique.
thumbnail Fig. 11

XRD pattern of the as-built sample.

Diffractogramme de l’échantillon brut de fabrication.

3.3 Microstructure and composition of the samples after heat treatment and HIP

Heat treatment of the as-built samples significantly modified their microstructures and crystallographic phases. In particular, cracks were no longer visible (Fig. 12a), because high temperatures promoted material diffusion, allowing for crack closure and sintering of contact surfaces. Carbide grain growth is measured between as-built and heat-treated sample. A mean WC grain size of 8.6 μm is measured for both heat-treated conditions (Figs. 12a and 12b).

However, heat treatment at ambient pressure had a limited effect on the porosity. According to the image analysis results, the relative density of the HT samples was 98.63%, with the disappearance of many of the smallest pores. In contrast to heat treatment at normal pressure the effect of HIP on porosity was remarkable. The microstructure shown in Figure 12b has a mass density of approximately 99.99%. Figure 12 presented before, shows the effect of post-treatment on the pore size and number for all analyzed cases. The apparent density results obtained using the Archimedes method are shown in Figure 13. Four samples were independently analysed for each condition. All measurements were conducted using acetone taking into account the temperature of the fluid. The highest density values were obtained for the HIP-treated samples, confirming the high efficiency of this process for pore closure. Nevertheless, the high relative density of 99.99% can also be related the loss of Co, which might evaporate during L-PBF.

The XRD patterns of the HT and HIP samples presented in Figure 14 are very similar. The phase transformation induced by high-temperature annealing and HIP resulted in the complete dissolution of the fragile W2C phase, which also has been found in previous studies. Meanwhile, formation of the Co6W6C phase with a cubic structure (Fd-3m) was observed in both HT and HIP samples. It can be concluded that the high pressure during HIP treatment did not significantly modify the phase transformation process and that the temperature (1450 °C) during treatments was the key factor influencing the final phase composition.

thumbnail Fig. 12

Sample microstructures after heat treatment at (a) ambient pressure and (b) HIP.

Microstructures des échantillons après traitement thermique à (a) pression ambiante et (b) HIP.
thumbnail Fig. 13

Mass density results obtained by the Archimedes method for the as-built, HT, and HIP-treated samples.

Résultats de la densité volumique obtenus par la méthode d’Archimède pour les échantillons brut de fabrication, traité à pression ambiante (HT) et traité par HIP.
thumbnail Fig. 14

XRD patterns of the (a) as-built, (b) HT and (c) HIP-treated samples.

Diffractogramme des échantillons (a) bruts de fabrication, (b) HT et (c) traités par HIP.

3.4 Macrohardness Vickers

Macrohardness tests were performed on the as-built, HT, and HIP-treated samples. Three measurements per sample were performed according to the standard methods, and the results are shown in Table 4. The mean value obtained for the as-built samples was 1096 HV30, which is close to those obtained for WC-17Co sintered reference parts [11,36]. In addition, material cracking was observed near the corners of the indent made on the sample cross-section, which indicated high material brittleness (Fig. 15a). The hardness of the HIP-treated (1152 HV30) samples was slightly higher than that of the as-built samples, likely due to their decreased porosity and material integrity improvement. In addition, material cracking was significantly reduced during indentation of the HT and HIP-treated specimens (Figs. 15b and 15c). This result can be related to the absence of the brittle W2C phase.

The standard deviation of the macrohardness for the HIP and HT samples were lower than those of the as-built samples due to their homogenized structures.

The results of the hardness test and the bulk density of the as-built and heat-treated samples were comparable to the sintered cermet (Tab. 4). This confirmed that the quality of the material produced by L-PBF is sufficiently high to provide suitable applications [11].

Table 4

Results of mechanical test.

Résultats des essais mécaniques.

thumbnail Fig. 15

Overview Vickers indent on the surface of the (a) as-built, (b) HT, and (c) HIP-treated samples.

Vue d’ensemble de l’empreinte Vickers sur la surface des échantillons (a) bruts de fabrication, (b) HT et (c) traités HIP.

4 Conclusion

The feasibility of using L-PBF additive manufacturing without a preheated plate to generate dense cermet samples from WC-17Co was successfully assessed. The results can be summarized as follows:

  • the manufacture of dense WC-17Co parts using L-PBF without preheated plate is possible;

  • samples with the lowest porosity (1.41%) and highest integrity were obtained using a volume energy density of 602 J/mm3 and a laser beam power of 130 W. The as-built samples contained the brittle W2C phase and some small-scale pores and cracks;

  • the sample structure, density, and hardness can be significantly improved through capsule-free HIP, resulting in close-to-zero porosity. Heat treatment performed at normal pressure eliminated short cracks and the small-scale porosity and enhanced the material hardness;

  • both post-treatment types (at atmospheric pressure and HIP) allowed the dissolution of the undesirable brittle W2C phase. Other complex carbides, such as Co3W3C, Co3W9C4, and Co6W6C, observed in the as-built samples were also present after simple heat treatment or HIP;

  • the macrohardness of the L-PBF-manufactured WC-17Co parts are very close to the values characterizing sintered reference samples;

  • the L-PBF parameters defined in this study can potentially be applied for the manufacture of complex-shaped parts of acceptable density and integrity without cracking.

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Cite this article as: Kevin Papy, Julien Favre, Alexey Sova, Andras Borbely, Philippe Bertrand, Jean-Marc Staerck, Manufacturing of WC-17Co 3D-objects by Laser Powder Bed Fusion followed by heat-treatment, Matériaux & Techniques 111, 304 (2023)

All Tables

Table 1

Chemical analysis of the initial powder according to ICP-MS.

Analyse chimique de la poudre initiale par procédé ICP-MS.

In the text
Table 2

Summary of powder properties.

Résumé des propriétés de la poudre.

In the text
Table 3

Ranges of L-PBF manufacturing parameters applied in the experiments.

Plages des paramètres du procédé L-PBF appliqués pour les optimisations paramétriques.

In the text
Table 4

Results of mechanical test.

Résultats des essais mécaniques.

In the text

All Figures

thumbnail Fig. 1

Global scheme of the study with three main steps: L-PBF process, post heat-treatment and hardness tests.

Schéma descriptif de l’étude en trois étapes principales : procédé L-PBF, post-traitement thermique et essais de dureté.
In the text
thumbnail Fig. 2

Intensity distribution surfaces at (a) 130 W and at (b) 300 W (laser power in the focal plane).

Distribution de l’intensité laser à (a) 130 W et à (b) 300 W (au plan focal).
In the text
thumbnail Fig. 3

SEM images of the WC-17Co particles at different magnification (a) x400 and (b) x3000.

Images MEB de la poudre de WC-17Co suivant différents grossissement (a) x400 et (b) x3000.
In the text
thumbnail Fig. 4

XRD pattern of the WC-17Co powder used for L-PBF process.

Diffractogramme de la poudre de WC-17Co utilisée pour le processus L-PBF.
In the text
thumbnail Fig. 5

Schematic illustration of (a) zigzag and (b) hexagonal scanning strategies.

Illustration schématique des stratégies de balayage (a) en aller/retour et (b) hexagonale.
In the text
thumbnail Fig. 6

Top: SEM images of samples manufactured via the hexagonal scanning strategy at (a) low, (b) medium, and (c) high energy densities. Bottom: Plot of sample density vs. energy density for zigzag and hexagonal scanning strategies.

En haut : images microscopes d’échantillons fabriqués selon la stratégie de balayage hexagonale à (a) faible, (b) moyenne et (c) haute densité d’énergie. En bas : évolution de la densité de l’échantillon en fonction de la densité d’énergie volumique pour les stratégies de balayage en aller/retour et hexagonale.
In the text
thumbnail Fig. 7

Image of the building plate and the samples manufactured for various analyses.

Image d’un plateau d’échantillons fabriqués pour diverses analyses.
In the text
thumbnail Fig. 8

SEM images of the as-built sample at (a) x500, (b) x1000, and (c) x5000 magnifications.

Images MEB de l’échantillon brut de fabrication aux grossissements (a) x500, (b) x1000, et (c) x5000.
In the text
thumbnail Fig. 9

Secondary electron SEM image of the as-built sample illustrating the presence of cracks, pores, and fine and coarse carbide grains.

Image MEB de l’échantillon brut de fabrication, illustrant la présence de fissures, de pores et de différentes tailles de grains de carbure.
In the text
thumbnail Fig. 10

Pore size distribution in the as-built samples and after heat treatment.

Distribution de la taille des pores dans les échantillons brut de fabrication et les échantillons après traitements thermique.
In the text
thumbnail Fig. 11

XRD pattern of the as-built sample.

Diffractogramme de l’échantillon brut de fabrication.
In the text
thumbnail Fig. 12

Sample microstructures after heat treatment at (a) ambient pressure and (b) HIP.

Microstructures des échantillons après traitement thermique à (a) pression ambiante et (b) HIP.
In the text
thumbnail Fig. 13

Mass density results obtained by the Archimedes method for the as-built, HT, and HIP-treated samples.

Résultats de la densité volumique obtenus par la méthode d’Archimède pour les échantillons brut de fabrication, traité à pression ambiante (HT) et traité par HIP.
In the text
thumbnail Fig. 14

XRD patterns of the (a) as-built, (b) HT and (c) HIP-treated samples.

Diffractogramme des échantillons (a) bruts de fabrication, (b) HT et (c) traités par HIP.
In the text
thumbnail Fig. 15

Overview Vickers indent on the surface of the (a) as-built, (b) HT, and (c) HIP-treated samples.

Vue d’ensemble de l’empreinte Vickers sur la surface des échantillons (a) bruts de fabrication, (b) HT et (c) traités HIP.
In the text

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