**1. Introduction**

The manufacturing and processing of materials is the basis of the modern economy and of civilization progress. Striving for sustainable development as well as limiting the harmful impact on the environment, including the reduction in CO<sup>2</sup> emissions and the reduction in energy consumption, leads to a continuous increase in the effectiveness and efficiency of technological processes, machines, devices, vehicles, tools, and building structures [1–3].

Therefore, new materials are being developed with higher mechanical properties, higher wear resistance and thus higher operating parameters. The conventional welding methods and technologies usually do not provide satisfactory results in applications such as joining and cladding of modern materials, e.g., high strength steel, nonferrous metals and alloys, composite materials, and nanostructured and hybrid materials [1,4–8].

The main disadvantages of typical methods of welding and cladding such as conventional arc or plasma arc methods are excessive heat input and the relatively large volume of the molten pool. This leads to overheating of the material, internal stresses and deformations, and unfavorable grain growth or dissolution of the reinforcing phase particles, e.g., carbides or nitrides. Therefore, the attention of researchers and industry is focused on the search for methods of joining and processing materials that ensure a minimum thermal effect on the material and the ability to precisely control the amount of heat and the thermal cycle of the process [9–13].

One of the dynamically developed areas of material processing technology in recent years is laser material processing technology. In the area of joining materials, laser welding shows several advantages because it provides high-power densities and low diameter of the laser beam spot, and thus high penetration depth, high welding speed, and low, precise, and controllable heat input. The above features are particularly important for the joining of modern and advanced alloys such as ultra-high-strength steels (UHSS), advanced high-strength steels (AHSS), modern stainless duplex and super duplex steel, nonferrous metals, and light metals, e.g., titanium and aluminum alloys [6,11,13–17].

Another area of laser beam application that is currently being intensely developed is the shaping of the properties of surface layers and the production of coatings for enhanced wear characteristics such as corrosion resistance, tribological properties, abrasive resistance, thermal and mechanical fatigue resistance, and resistance to impact load [1–8,18].

In the area of surface treatment and cladding, the laser beam as a heat source is also advantageous. The most significant advantages include low and controllable heat input, and thus limited thermal impact on the substrate, reduced internal stresses, minimized distortions, high solidification and cooling rates, low penetration depth, and low dilution of the clad layer by the substrate material. Low heat inputs and related high cooling rates during laser processing are beneficial in many applications of surface treatment and cladding as they provide high quality clads and surface layers. Usually, surface layers are characterized by superior metallurgical bonding and a fine-grained and refined microstructure, which are decisive for providing superior tribological characteristics, high abrasive or erosion wear resistance, and corrosion or fatigue resistance [1,8,18].

It is worth noting that laser-made clads and surface layers show higher functional properties compared to surface layers made by plasma arc (PTA) or conventional arc cladding processes, even if the same additional materials are applied. This is due to intense grain refinement caused by very fast cooling and solidification of the molten pool, typically an order of magnitude higher than plasma (PTA) and conventional arc cladding (e.g., GTA–gas tungsten arc) [1,12].

Thanks to the advantages mentioned above, lasers are used in different processes of surface treatment such as laser surface hardening (LSH), melting (LSM), shocking (LSS), texturing (LST), alloying (LSA), cladding (LSC), remelting (LSR), and surface deposition (LSD), and in additive manufacturing processes such laser metal deposition (LMD) or selective laser sintering and melting (SLS/SLM) [1].

Continuous development of laser generators and related improvement of laser radiation and laser beam characteristics enable further development in the field of laser material processing, including laser welding and laser surface treatment. In addition to the gaseous CO<sup>2</sup> and solid-state rod type Nd:YAG laser generators used for the last few decades, new generations of solid-state lasers are currently available such as high-power diode lasers (HPDL), disk lasers, and fiber lasers [1,14–17]. Modern generations of laser devices offer new or expanded technological possibilities. For example, as shown in the literature, high-power diode lasers with a square or rectangular beam spot and uniform energy distribution (multimode beam) are favorable for surface treatment applications. The design of diode lasers has been developed to the extent that today there are direct diode lasers (DDL) which emit single-mode radiation with power up to 8 kW, suitable for cutting applications. Moreover, the single-mode DDL lasers are even competitive with CO<sup>2</sup> gas lasers and disk lasers, as well as fiber lasers [1,14,17,18].

There has also been significant progress in the design and construction of fiber generators in the last decade. Modern fiber lasers are also beneficial due to their compact design, high wall plug efficiency, low beam divergence, and possibility to focus the beam into a very small spot, even at relatively high power [18]. Moreover, these fiber lasers can provide a single-mode beam up to 20 kW, while the multimode industrial applications reach a power level of 100 kW. As a result, lasers are currently used for one-sided or two-sided welding, even for thick steel plates with a thickness up to approx. 30 mm.

Another interesting example of dynamic development in the field of material processing technology is hybrid processes. According to the definition, the hybrid process combines at least two different machining processes which must be carried out simultaneously and in the same processing area (e.g., one weld pool created during fusion welding). In the case of welding, the laser beam is often combined with the arc, e.g., laser + GMA (gas metal arc), laser + GTA (gas tungsten arc) or laser + PTA (plasma tungsten arc). Such combinations of two different heat sources allow taking the advantages of both welding methods. The laser beam is responsible for deep penetration of the material, while the arc processes provide additional heat and material to fill the gap. Consequently, it is possible to weld at high speed and ensure the correct shape of the weld at high depth. Hybrid processes are also being developed in the field of surface treatment and coating applications, ensuring higher functional parameters or the possibility of producing coatings with special properties [1,18].

#### **2. Short Characterization of the Special Issue**

The Materials and Methods should be described with sufficient details to allow others to replicate and build on the published results. Please note that the publication of your manuscript implicates that you must make all materials, data, computer code, and protocols associated with the publication available to readers. Please disclose at the submission stage any restrictions on the availability of materials or information. New methods and protocols should be described in detail, while well-established methods can be briefly described and appropriately cited.

This Special Issue, entitled "Development of Laser Welding and Surface Treatment of Metals", is a complementary and valuable resource of knowledge in the following fields:


In the work presented by Wang et al. [19], plates of X100 pipeline steel with a thickness of 12.8 mm were welded by a high-power laser on a robotized cell. The authors provided a quantitative correlation between thermal cycling and the microstructure of the welded joint by means of empirical and numerical study. They also showed that the effects of austenitization are more significant than those of the cooling rate on the final microstructures of the laser-welded joint of the investigated steel grade. The results of the presented work can provide scientific guidance and a reference for the simulation of the temperature field during laser welding, especially in the case of laser welding of X100 pipeline steel grade.

Kik [20] presented different computational techniques for comparison in the case of single and multipass arc and laser beam welding processes. The numerical analyses were conducted by the SYSWELD software package and the author showed differences between the applied computational techniques, as well as the benefits and disadvantages of the applied computational techniques. The presented results of the numerical simulations may be used for preparing an efficient and fast optimization of the laser and related fusion welding processes, according to the criteria of minimizing deformations or identification of potential defects of structures.

In the second article, Kik [23] presented new possibilities for modifying heat source models in numerical simulations of laser welding processes conducted using VisualWeld (SYSWELD) software. He proposed a modification of heat source models and methods of defining the heat introduced into welded materials using solid-state and high-power diode lasers. He provided calibration and validation of the proposed models of heat sources based on metallographic tests and thermal cycles; therefore, the presented results may be helpful for further numerical analysis of laser processing by different beam shape and energy distribution.

The study presented by Górka [21] was focused on investigation of the influence of the basic parameters of autogenous laser welding of 700MC 10.0 mm thick plates on their microstructure and mechanical properties. He provided a comprehensive analysis of the microstructure and a detailed analysis of the phase constituents. The quality of the test joint was estimated as meeting level B, in accordance with 13919-1 standard. He concluded that in cases of S700MC steel, the analysis of the phase transformation of austenite exposed to welding thermal cycles and the value of carbon equivalent cannot be the only factors taken into consideration when assessing weldability.

The laser welding process was also investigated by Danielewski and Skrzypczyk [22]. They presented results of numerical and experimental analyses of the gaseous CO<sup>2</sup> laser welding of lap joints of low-carbon structural steel. They determined hardness profiles and weld geometry for the welds produced at different parameters. They also provided detailed analysis of the metallographic structure of fusion zone (FZ) and heat-affected zone (HAZ), as well as a quality assessment of the tested joints.

The study presented by Landowski et al. [24] concerned fiber laser welding of dissimilar butt joints between 316 L austenitic and 2304 lean duplex stainless steel plates of 8.0 mm thickness. Based on nondestructive tests and metallographic examinations, they showed that the tested joints meet the acceptance criteria for level B, in accordance with EN ISO 13919–1 standard. Finally, they proved that the elaborated procedure can be applied for welding of 316 L–2304 stainless steel plates with a thickness of 8.0 mm without ceramic backing. The results also provide valuable guidance for further research in this field.

The results of the study presented by Pa´ncikiewicz et al. [25] also concerned laser welding of dissimilar joints of AISI 430F (X12CrMoS17) martensitic stainless steel and AISI 304 (X5CrNi18-10) austenitic stainless steel. They provided detailed analysis of the microstructure of the base metals and the dissimilar test joint. However, the test joint meets the quality level C, in accordance with the ISO 13919-1 standard, due to gas pores in the weld metal.

Another study on laser welding was presented by Tofil et al. [26]. They conducted a test of laser welding of thin-walled AISI 316 L steel pipes, with diameters of 1.5 and 2.0 mm, used in medical equipment. They determined the range of optimal parameters of laser welding with a SISMA LM-D210 Nd:YAG laser. The results of the study showed good joint properties with a strength of more than 75% in the thinner pipe, uniform distribution of alloying elements, and a complex dendritic structure characteristic of pulse welding. Therefore, the presented results may also be guidance for further research in this field.

The topic of laser surface treatment and cladding begins with Lisiecki's article [27] on the study of the optical properties of surface layers produced on the titanium alloy substrate by surface melting and nitriding. The aim of this study was to determine the optical properties of the surface of titanium alloy Ti6Al4V. The specular, diffuse, and total reflection was determined for different surface conditions (roughness and oxides layer) and after laser surface treatment. The experiments were conducted for 808 nm wavelength, typical for diode lasers. The main conclusion of the study was that the distinct differences in absorption affected the heat transfer and thermal conditions of laser heating, and thereby the penetration depth during laser melting and nitriding of the titanium alloy.

A highly interesting study on laser surface modification of an aluminum alloy by boron carbides was carried out by Sroka et al. [28]. The fiber laser YLS-4000 with a maximum output power of 4.0 kW was used in this study. The effect of basic laser-alloying parameters of the foundry AlMg9 alloy with the B4C particle on the microstructure and chosen properties was investigated and determined.

Another interesting study on laser cladding was performed by Li et al. [29]. In their study, the high aluminum and chromium Fe-B-C coating was produced by laser cladding on a 2Cr13 steel substrate. They investigated the microstructure and tribological characteristics. The results showed that the coating exhibited excellent wear resistance compared to the substrate. They concluded that the results could provide technical support in improving the properties of the Fe-based laser-cladded coating.

The next study, presented by Majkowska-Marzec et al. [30], is related to investigation of the mechanical and corrosion properties of a laser surface-treated titanium alloy with carbon nanotube coating. The coating was initially formed by electrophoretic deposition (EPD) and then subjected to laser treatment. The authors pointed to some advantages of such a coating. The main advantages were increased corrosion resistance, a lack of surface cracks, increased strength, and favorable contact angles between 46◦ and 82◦ (resulting in hydrophilic surfaces suitable for cell adhesion).

A comparative study of laser and plasma cladding was presented by Czupry´nski et al. [31]. The coatings of Inconel 625 on the 16Mo3 steel pipes were produced by plasma (PTA) cladding and high-power direct diode laser (HPDDL) cladding. The authors determined the quality of the coating by NDT testing, metallographic examinations, and microstructure analysis. They concluded that both tested and compared methods can provide high-quality protective coatings that may be operated at elevated temperatures up to 625 ◦C.

Another study on the effect of laser treatment on tribological properties was presented by Lubas et al. [32]. They analyzed the influence of different engine oils on the tribological parameters of sliding couples with a laser-borided surface layer produced on specimens of AISI 5045 steel, by laser remelting of a surface layer coated with amorphous boron. The tribological and wear characteristics were investigated on the pairs of AISI 5045 steel and SAE-48 bearing alloy, lubricated with 5W-40 and 15W-40-type engine oils. The detailed analysis of wear mechanisms was identified and described.

The thematic series on laser surfacing and cladding ends with the article presented by Lisiecki and Slizak [ ´ 33] on hybrid laser deposition of composite coatings at cryogenic conditions. This study demonstrated the effect of forced and localized cooling by a nitrogen vapor stream under cryogenic conditions during laser deposition of WC-Ni powder on the geometry and microstructure of clad layers and the dry sliding wear resistance of the coatings. Additionally, the quantitative relationship between heat input, cooling conditions, and carbide grain size distribution, as well as carbide share in relation to the matrix, was determined. It was shown that the novel demonstrated technique of localized forced cooling during laser cladding has advantageous effects. The forced localized cooling provides approximately 20% lower penetration depth and dilution, decreases tendency for tungsten carbide decomposition, and provides more uniform distribution and a higher share of massive eutectic W2C-WC carbides across the coating.

Another very novel and interesting subject related to arc welding processes was presented by Szczucka-Lasota and Szymczak et al. [34,35]. The study presented in two articles was related to the application of a novel and patented technique of localized cooling of the weld by micro-jet streams. The first presented article is focused on the influence of processing conditions, especially localized cooling on the properties of the weld of Docol 1200 M steel. They found that the novel technique of cooling can provide a high fatigue limit at the level of at least 480 MPa. The second article presented by Szymczak et al. [35] is related to a similar issue but a different type of steel: S960MC sheet of 2.0 mm thickness. In this case, the Wöhler S–N curve of the weld was determined, indicating that the value of the fatigue limit of the tested weld was 100 MPa. The weld at the Union NiMoCr welding wire was indicated as the joint with the highest resistance on static and fatigue loadings.

Some manuscripts dealt with other arc processes and technologies, but the presented solutions were very innovative. An example of a highly interesting subject was presented by Tomków et al. [36]. They investigated underwater welding in terms of the role of the bead sequence during welding under such conditions. The test beads were produced by covered electrodes with the use of normalized S355G10 + N steel. They reported that welding in the underwater environment carries many problems related to arc stability, and thus the properties of the welds. Moreover, they revealed the effects of refining and tempering the microstructure in heat-affected zones of the previous beads, which resulted in a reduction in hardness.

The study presented by Mician et al. [37] was related to the influence of bead sequence on the deposit and heat-affected zone properties during arc surfacing. The process applied to surfacing was gas metal arc welding (GMAW) and the substrate plate was an S355 steel plate. The authors provided a comparative analysis of structure and hardness, considering the thermal impact of the bead sequence. As a result of the study, the most favorable sequence in terms of structure and hardness distribution, maximum hardness, and range of hardness has been established.

The second study presented by Czupry´nski [38] was focused on flame spraying of aluminum coatings reinforced by carbon nanotubes as an alternative for laser cladding technology. The microstructure of the coatings was analyzed, as well as abrasion and erosion resistance. He concluded that the flame spraying of aluminum coatings reinforced with the carbon nanotubes can be an effective alternative for laser cladding technology. Moreover, such coatings can be implemented in the automotive industry for the production of components characterized by high strength, wear resistance, good thermal conductivity, and low density.

The next article provided by Czupry ´nski [39] presented results of the investigations of abrasive wear characteristics of wear-resistant plates made by cladding with innovative tubular electrodes with a metallic core and an experimental chemical composition. Detailed analyses of microstructure and wear tests of different reference plates were conducted. The author showed a high metal–mineral abrasive wear resistance of the deposit produced by the experimental tubular electrode.

Another original issue was presented in the study by Kapustynskyj et al. [40]. They focused on the numerical and analytical study of the effects of laser surface processing of thin-plate low-carbon structural steel. The authors developed the original analytical methodology of the estimation of the cross-section area of the laser-processed metal, that can be applied to the choosing of a reasonable distance between the centers of the laser-processed tracks. The results of the experimental and finite element numerical and analytical analyses showed that the laser treatments of the surface of the steel plate increase the yield point of the laser-processed region, as well as the axial and flexural stiffness of the plate.

#### **3. Concluding Remarks**

This Special Issue was very successful due to the valuable articles that were submitted, the wide variety of topics and research problems that were undertaken, and in-depth analysis of the state of the art in different fields of laser and related processes. This Special Issue contains a significant number of research articles (22 in total); however, the entire number of manuscripts submitted to this Special Issue was half that higher. Unfortunately, some of the manuscripts have not gone through a very rigorous review process. Such a large interest and quantity of articles shows the importance of the issue and themes. It should also be emphasized that each research article in the field of materials processing is the result of tedious, long-term, and interdisciplinary research, usually conducted by a team of scientists.

The topics of the articles were mainly focused on laser welding, laser cladding, and laser surface treatment, as well as related processes of fusion welding and manufacturing of coatings.

The wide range of experimental, numerical, or analytical study conducted and presented by 58 authors, representing 19 academic or research centers and 3 industrial centers from 5 countries (Poland, China, Slovakia, Lithuania, and Ukraine) and 16 cities, proved the topicality and importance of this Special Issue.

**Funding:** This publication received no external funding.

**Acknowledgments:** As the Editor of the Special Issue, I would like to thank all the authors of the submitted articles, as well as the reviewers, editors, and everyone who contributed to publishing the Special Issue.

**Conflicts of Interest:** The author declares no conflict of interest.

### **References**


**Aleksander Lisiecki 1,\* and Dawid Slizak ´ <sup>2</sup>**


**Abstract:** The purpose of this study was to demonstrate the effect of forced and localized cooling by nitrogen vapours stream under cryogenic conditions during laser deposition of WC-Ni powder on the geometry, microstructure of clad layers and dry sliding wear resistance of the coatings. For this purpose, comparative tests were performed by conventional laser cladding at free cooling conditions in ambient air and by the developed novel process of laser deposition with additional localized cooling of the solidifying deposit by nitrogen vapours stream. Due to presence of gaseous nitrogen in the region of the melt pool and solidifying deposit, the process was considered as combining laser cladding and laser gas nitriding (performed simultaneously), thus the hybrid process. The influence of the heat input and cooling conditions on the geometrical features, dilution rate, share of carbides relative to the matrix, and the fraction share of carbides, as well as hardness profiles on cross sections of single stringer beads was analysed and presented. The XRD, EDS analysis and the sieve test of the experimental powder were used to characterize the composite WC-Ni type powder. The OM, SEM, EDS and XRD test methods were used to study the microstructure, chemical and phase composition of clad layers. Additionally, ball-on-disc tests were performed to determine the wear resistance of representative coatings under dry sliding conditions. The results indicate that the novel demonstrated technique of localized forced cooling of the solidifying deposit has advantageous effect, because it provides approximately 20% lower penetration depth and dilution, decreases tendency for tungsten carbides decomposition, provides more uniform distribution and higher share of massive eutectic W2C-WC carbides across the coating. While the conventionally laser cladded layers show tendency for decomposition of carbide particles and resolidifying dendritic complex carbides mainly M2C, M3C and M7C<sup>3</sup> containing iron, nickel, and tungsten, and with Ni/Ni3B matrix. The quantitative relationship between heat input, cooling conditions and the carbides grain size distribution as well as carbides share in relation to the matrix was determined.

**Keywords:** laser cladding; laser deposition; hybrid process; cryogenic conditions; composite coatings; WC-Ni coatings; fiber laser

#### **1. Introduction**

Laser cladding offers some advantages over the other methods of cladding or coatings [1–10]. The most significant advantages are high power density, localized and precise heating, high scanning speed, thus low heat input, low penetration depth and low dilution. Laser cladding is often used to produce wear-resistant layers based on metallic and composite materials. One group of composite materials are metal matrix composites (MMC) based on nickel matrix reinforced with tungsten carbides. Such composites are characterized by good wear resistance, corrosion resistance and also satisfactory dynamic load resistance. The high abrasive wear resistance is provided by hard particles of tungsten carbides distributed in the ductile nickel matrix. In turn, the ductile matrix provides the

**Citation:** Lisiecki, A.; Slizak, D. ´ Hybrid Laser Deposition of Composite WC-Ni Layers with Forced Local Cryogenic Cooling. *Materials* **2021**, *14*, 4312. https:// doi.org/10.3390/ma14154312

Academic Editor: Antonio Santagata

Received: 31 May 2021 Accepted: 26 July 2021 Published: 2 August 2021

**Publisher's Note:** MDPI stays neutral with regard to jurisdictional claims in published maps and institutional affiliations.

**Copyright:** © 2021 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https:// creativecommons.org/licenses/by/ 4.0/).

ability to withstand high loads, especially compressive stresses. Moreover, such composites show satisfactory cohesion between hard WC phases and the matrix of Ni solution, due to good wettability of WC particles by the Ni-based alloys. The process of laser cladding is considered as advantageous for manufacturing the composite clad layers due to low heat input, thus limited reaction between the ceramic phases and the metallic matrix [11–14]. However, even in the case of laser cladding, especially in the output power range of several kilowatts, the heat input may be too high producing unfavourable microstructure due to partial and even complete melting of WC particles. Moreover, under specific conditions the WC particles tend to fall off in the melt pool, causes uneven distribution of the carbides of various size across the surface layer. Zhang et al. indicate that the accumulation of WC particles at the bottom of the clad coating is detrimental for the properties of the coating [12]. They also point that it is the biggest obstacle to engineering applications of such composite coatings. Various methods for providing the uniform distribution of ceramic particles in the composite coatings, such mechanical vibration, ultrasonic, magnetic fields were studied and described in the literature [13]. Li et al. demonstrated an original technique of laser cladding of WC-Ni composite coatings assisted by micro-vibrations generated by a vibration exciter system made of magneto strictive materials [14,15]. They pointed the beneficial effect of the micro-vibrations on microstructure, including uniform distribution of tungsten carbides [15]. In turn Huang et al. demonstrated a cladding technique involving a pulsed Nd:YAG laser for reducing the heat input [16]. They successfully produced dense and crack-free WC-Ni composite clad layers on H13 substrate with the thickness up to 1.0 mm.

Another technique that allows to reduce the effect of the heat on the material and to control the solidification rate is a forced cooling. An original technique of micro-jet cooling by compressed gas streams of the deposit during arc cladding was elaborated and demonstrated by W˛egrzyn et al. [17]. They proved the beneficial effect of the forced cooling of the deposit at solidification stage resulted in refinement of microstructure and enhanced wear resistance.

Liquid nitrogen bath providing cryogenic conditions of cooling was also tested by some researchers during laser surface melting or heat treatment, mainly for nonferrous alloys [18–23]. Such technique of liquid nitrogen bath cooling of the substrate was adopted and developed by the authors for laser powder deposition of metallic and composite coatings, as demonstrated in several previous publications [1,24–31]. It is worth to note, that complexity and difficulty of the laser powder deposition under such cryogenic conditions with coaxial powder delivery into the melt pool is significantly higher than the laser surface melting of the substrate supercooled by liquid nitrogen bath. Moreover, the results obtained so far indicate that due to intensive evaporating of liquid nitrogen, the gaseous nitrogen present in the region of powder deposition and melt pool provides an active atmosphere typical for laser gas nitriding (LGN) or laser alloying (LA). Therefore, the proposed new method of laser coating is considered as a hybrid process combining laser powder deposition and laser gas nitriding, additionally conducted under cryogenic conditions.

The preliminary tests of laser cladding of composite WC-Ni coatings, conducted by the authors, have shown that the forced cooling of the substrate by liquid nitrogen bath can provide some beneficial effects. The most significant are limited penetration depth, limited dilution, limited tendency for tungsten carbide melting, uniform and dense distribution of carbides, favourable hardness distribution and high values of hardness. However, some disadvantages and limitations of the liquid nitrogen bath (volume cooling) applied for laser powder deposition were identified. Due to high extent of supercooling the substrate and intensive dissipation of the heat during laser heating, the tendency for incomplete or lack of penetration of the clad was observed, especially at the lowest heat input range. Therefore, another technique of forced cooling by means of liquid nitrogen based on stream local cooling of the deposition region, that allows controlled and less intensive cooling of the substrate was developed and investigated by the authors.

According to the authors' knowledge and experience, the presented results are fully original and unique, and the results of similar studies have not been available nor published so far. According to the authors' knowledge and experience, the presented results are fully original and unique, and the results of similar studies have not been available nor published so far.

#### **2. Materials and Methods 2. Materials and Methods**

#### *2.1. Materials*

The non-alloy structural steel S235JR (according to EN 10025-2) was chosen as the substrate for technological tests. This grade of steel with a low content of alloying elements, good plasticity and excellent weldability was chosen to minimize the influence of the substrate material on the cladded layer, especially minimize the internal stresses and tendency for cracking of clad layers, Table 1. Specimens for laser cladding tests were cut from a steel plate 5.0 mm thick into coupons with dimensions 100 <sup>×</sup> 100 mm<sup>2</sup> . *2.1. Materials*  The non-alloy structural steel S235JR (according to EN 10025-2) was chosen as the substrate for technological tests. This grade of steel with a low content of alloying elements, good plasticity and excellent weldability was chosen to minimize the influence of the substrate material on the cladded layer, especially minimize the internal stresses and tendency for cracking of clad layers, Table 1. Specimens for laser cladding tests were cut

**Table 1.** Chemical composition of non-alloy structural steel S235JR (EN 10025-2) based on the supplier's (Thyssenkrupp, D ˛abrowa Górnicza, Poland)) certificate (wt %). **Table 1.** Chemical composition of non-alloy structural steel S235JR (EN 10025-2) based on the sup-

from a steel plate 5.0 mm thick into coupons with dimensions 100 × 100 mm2.


In turn, the powder for cladding was experimentally composed in such a way as to ensure high abrasion resistance, and corrosion resistance of deposit even at elevated temperatures. Therefore, tungsten carbides, as the reinforcing phase, were mixed with nickel-based powder providing the ductile metal matrix characterized by high resistance for corrosion and high temperature. The nickel-based metal powder was gas atomized in argon atmosphere and the spherical particles size range was 63 ÷ 160 µm. While the tungsten carbide irregular shaped (crushed) particles size range was 100 ÷ 160 µm. The sieve test of the mixed experimental powder was performed for statistical analysis according to the standard PN-EN 24497/ISO 4497. Sieve shaker LPzE-2e (Multiserw, Brze´znica, Poland) and the moisture analyser with the laboratory scales BTS110 (AXIS, Gda ´nsk, Poland) were applied in the sieve test. The sieve test was repeated three times for different batches of powder weighing approx. 100 g each, for statistical reasons. The mass fraction share and cumulative particle size distribution curve determined for the experimentally composed powder are presented in Figure 1. The morphology, chemical and phase composition of the experimental powder are presented in Figures 2 and 3. In turn, the powder for cladding was experimentally composed in such a way as to ensure high abrasion resistance, and corrosion resistance of deposit even at elevated temperatures. Therefore, tungsten carbides, as the reinforcing phase, were mixed with nickelbased powder providing the ductile metal matrix characterized by high resistance for corrosion and high temperature. The nickel-based metal powder was gas atomized in argon atmosphere and the spherical particles size range was 63 ÷ 160 µm. While the tungsten carbide irregular shaped (crushed) particles size range was 100 ÷ 160 µm. The sieve test of the mixed experimental powder was performed for statistical analysis according to the standard PN-EN 24497/ISO 4497. Sieve shaker LPzE-2e (Multiserw, Brzeźnica, Poland) and the moisture analyser with the laboratory scales BTS110 (AXIS, Gdańsk, Poland) were applied in the sieve test. The sieve test was repeated three times for different batches of powder weighing approx. 100 g each, for statistical reasons. The mass fraction share and cumulative particle size distribution curve determined for the experimentally composed powder are presented in Figure 1. The morphology, chemical and phase composition of the experimental powder are presented in Figures 2 and 3.

**Figure 1.** The mass fraction share (**a**) and cumulative particle size distribution curve (**b**) determined by sieve test (PN-EN 24497/ISO 4497) for the experimental composite WC-Ni powder, Table 2. **Figure 1.** The mass fraction share (**a**) and cumulative particle size distribution curve (**b**) determined by sieve test (PN-EN 24497/ISO 4497) for the experimental composite WC-Ni powder, Table 2.

nickel-based metal matrix (wt %), Figures 2 and 3.

**Table 2.** Nominal composition of experimental powder as a mixture of tungsten carbides with

**WC Fe Si B C Ni**  60 2.0 3.0 3.0 0.02 Bal.


**Table 2.** Nominal composition of experimental powder as a mixture of tungsten carbides with nickel-based metal matrix (wt %), Figures 2 and 3.

**Figure 2.** Morphology of the experimental composite WC-Ni powder; mixture of irregular WC particles and spherical metallic Ni-based particles, Table 2. **Figure 2.** Morphology of the experimental composite WC-Ni powder; mixture of irregular WC particles and spherical metallic Ni-based particles, Table 2. **Figure 2.** Morphology of the experimental composite WC-Ni powder; mixture of irregular WC particles and spherical metallic Ni-based particles, Table 2.

**Figure 3.** XRD pattern of the experimental WC-Ni composite powder, Table 2, Figures 1 and 2. **Figure 3.** XRD pattern of the experimental WC-Ni composite powder, Table 2, Figures 1 and 2. **Figure 3.** XRD pattern of the experimental WC-Ni composite powder, Table 2, Figures 1 and 2.

#### *2.2. Laser Deposition Tests 2.2. Laser Deposition Tests 2.2. Laser Deposition Tests*

The tests of laser deposition were performed by means of a prototype robotized stand. The stand was equipped with a six-axis robot Panasonic GII TL-190 (Industrial Solutions Company, Panasonic Corporation, Osaka, Japan) with maximum load capacity 6.0 kg, high power fiber laser (HPFL) IPG YLS-3000-CT-Y15 (IPG Photonics, Oxford, MS, USA) with maximum output power 3.0 kW, emitting at a wavelength 1.07 µm, and characterized by a single mode energy distribution across the spot (Gaussian intensity distribution), custom made powder delivery system, and a specially designed system for localized delivery of liquid nitrogen or nitrogen vapours stream. The laser beam was transmitted from the laser generator to the focussing head IPG FLW by 20 m long HLC-8 fiber with a core diameter 200 µm. The focal length of the applied FLW head was 200 mm, while the collimator focal length was 100 mm. The beam parameter product (BPP) of the laser The tests of laser deposition were performed by means of a prototype robotized stand. The stand was equipped with a six-axis robot Panasonic GII TL-190 (Industrial Solutions Company, Panasonic Corporation, Osaka, Japan) with maximum load capacity 6.0 kg, high power fiber laser (HPFL) IPG YLS-3000-CT-Y15 (IPG Photonics, Oxford, MS, USA) with maximum output power 3.0 kW, emitting at a wavelength 1.07 µm, and characterized by a single mode energy distribution across the spot (Gaussian intensity distribution), custom made powder delivery system, and a specially designed system for localized delivery of liquid nitrogen or nitrogen vapours stream. The laser beam was transmitted from the laser generator to the focussing head IPG FLW by 20 m long HLC-8 fiber with a core diameter 200 µm. The focal length of the applied FLW head was 200 mm, while the collimator focal length was 100 mm. The beam parameter product (BPP) of the laser The tests of laser deposition were performed by means of a prototype robotized stand. The stand was equipped with a six-axis robot Panasonic GII TL-190 (Industrial Solutions Company, Panasonic Corporation, Osaka, Japan) with maximum load capacity 6.0 kg, high power fiber laser (HPFL) IPG YLS-3000-CT-Y15 (IPG Photonics, Oxford, MS, USA) with maximum output power 3.0 kW, emitting at a wavelength 1.07 µm, and characterized by a single mode energy distribution across the spot (Gaussian intensity distribution), custom made powder delivery system, and a specially designed system for localized delivery of liquid nitrogen or nitrogen vapours stream. The laser beam was transmitted from the laser generator to the focussing head IPG FLW by 20 m long HLC-8 fiber with a core diameter 200 µm. The focal length of the applied FLW head was 200 mm, while the collimator focal length was 100 mm. The beam parameter product (BPP) of the laser beam determined

2Theta (°)

M23C6

M7C3

beam determined for the applied optics configuration was 5.8 mm⋅mrad. The BPP was measured by Primes Focus Monitor FM-120 (PRIMES GmbH, Pfungstadt, Germany).

nitrogen bath, to overcome the encountered limitations of such technique of volume cryogenic cooling [1]. During the previous preliminary tests, a wide range of processing parameters was investigated, at different scanning speed and output laser power. In this study the parameters were narrowed and optimized in order to provide the width of a

The stream localized cooling system was designed based on the experiences of pre-

beam determined for the applied optics configuration was 5.8 mm⋅mrad. The BPP was measured by Primes Focus Monitor FM-120 (PRIMES GmbH, Pfungstadt, Germany).

nitrogen bath, to overcome the encountered limitations of such technique of volume cryogenic cooling [1]. During the previous preliminary tests, a wide range of processing parameters was investigated, at different scanning speed and output laser power. In this study the parameters were narrowed and optimized in order to provide the width of a

The stream localized cooling system was designed based on the experiences of pre-

for the applied optics configuration was 5.8 mm·mrad. The BPP was measured by Primes Focus Monitor FM-120 (PRIMES GmbH, Pfungstadt, Germany).

The stream localized cooling system was designed based on the experiences of previous preliminary study of laser cladding the steel substrate immersed partially in liquid nitrogen bath, to overcome the encountered limitations of such technique of volume cryogenic cooling [1]. During the previous preliminary tests, a wide range of processing parameters was investigated, at different scanning speed and output laser power. In this study the parameters were narrowed and optimized in order to provide the width of a single bead approx. 2.0 ÷ 3.0 mm, and proper fusion to the steel substrate, Table 3, Figures 4 and 5. *Materials* **2021**, *14*, x FOR PEER REVIEW 5 of 26 single bead approx. 2.0 ÷ 3.0 mm, and proper fusion to the steel substrate, Table 3, Figures 4 and 5.

**Table 3.** Parameters of laser deposition of experimental composite WC-Ni powder at free cooling conditions (conventional laser cladding) and forced localized cooling by nitrogen vapours stream under cryogenic conditions (hybrid laser deposition process), Figure 6. **Table 3.** Parameters of laser deposition of experimental composite WC-Ni powder at free cooling conditions (conventional laser cladding) and forced localized cooling by nitrogen vapours stream under cryogenic conditions (hybrid laser deposition process), Figure 6.


Remarks: UB—uneven bead, SP—single pore, IF—incomplete fusion, V—voids, HQ—high quality, LF—lack of fusion. \* energy input is calculated by simply dividing the laser power by scanning speed, while the heat input should also include the heat transfer efficiency. Other process parameters; powder feeding rate: 8.5 g/min, diameter of the nozzle tip for nitrogen vapours stream: 5.0 mm, average consumption of liquid nitrogen: 80 ÷ 100 g/min. energy input is calculated by simply dividing the laser power by scanning speed, while the heat input should also include the heat transfer efficiency. Other process parameters; powder feeding rate: 8.5 g/min, diameter of the nozzle tip for nitrogen vapours stream: 5.0 mm, average consumption of liquid nitrogen: 80 ÷ 100 g/min.

**Figure 4.** (**a**) A view of the mount of three coaxial nozzles for powder delivery and injection into the melt pool on the laser focusing head; and (**b**) a view of the hybrid laser deposition of composite WC-Ni powder at cryogenic conditions by the localized cooling with the nitrogen vapours stream and the head with three coaxial nozzles. **Figure 4.** (**a**) A view of the mount of three coaxial nozzles for powder delivery and injection into the melt pool on the laser focusing head; and (**b**) a view of the hybrid laser deposition of composite WC-Ni powder at cryogenic conditions by the localized cooling with the nitrogen vapours stream and the head with three coaxial nozzles.

> (**a**) (**b**) Manometer and safety valve Liquid nitrogen tank Control panel Laser head Laser beam Steel substrate Powder stream Powder nozzle injection Liquid nitrogen stream Insulated hose Nozzle of liquid nitrogen Melt pool Clad Scanning direction Prior to the laser cladding tests, the specimens were sandblasted (surface roughness Ra 25 ÷ 60 µm) and degreased with acetone to remove surface contamination and provide stable and repeatable surface conditions for absorption of laser radiation. Surface roughness was determined by a portable surface roughness tester SJ-210 Surftest (Mitutoyo Corporation, Kanagawa, Japan). The test cladded layers were produced as single stringer beads by laser deposition of the experimentally composed powder at different heat input and different conditions of cooling. The length of the individual bead was 80 mm, while the shift between beads was 20 mm, Figure 6. One set of test clad layers was produced during conventional laser cladding with natural cooling (so called free cooling) of the specimen in the ambient air. The second set of test clad layers was produced at the same processing parameters but additionally with the application of forced cooling by localized

Prior to the laser cladding tests, the specimens were sandblasted (surface roughness Ra 25 ÷ 60 µm) and degreased with acetone to remove surface contamination and provide

**Figure 5.** (**a**) A view of the experimental setup for laser powder deposition of composite layers with forced local cryogenic cooling by the nitrogen vapours stream (from left: cryogenic dewar with automatic control unit, elastic hose for nitrogen

der deposition with forced cooling of the substrate by localized nitrogen vapours stream.

tion process), Figure 6.

**No. Surface Layer Indication Free Cooling/Hybrid** 

> stream of liquid nitrogen. The temperature of the liquid nitrogen in the container was below −196 ◦C, while the temperature of the nitrogen vapours stream measured at the tip of the nozzle by a thermocouple was approx. −160 ÷ 164 ◦C. In such a way deep cryogenic treatment (DCT) was provided in the localized area of forced cooling, Table 3, Figure 4b. (**a**) (**b**) **Figure 4.** (**a**) A view of the mount of three coaxial nozzles for powder delivery and injection into the melt pool on the laser focusing head; and (**b**) a view of the hybrid laser deposition of composite WC-Ni powder at cryogenic conditions by the localized cooling with the nitrogen vapours stream and the head with three coaxial nozzles.

*Materials* **2021**, *14*, x FOR PEER REVIEW 5 of 26

**Table 3.** Parameters of laser deposition of experimental composite WC-Ni powder at free cooling conditions (conventional laser cladding) and forced localized cooling by nitrogen vapours stream under cryogenic conditions (hybrid laser deposi-

1 LC1/HC1 500 500 60 SP/HQ 2 LC2/HC2 500 1000 120 SP, HQ/HQ 3 LC3/HC3 500 1500 180 SP, HQ/HQ 4 LC4/HC4 500 2000 240 SP/SP, HQ Remarks: UB—uneven bead, SP—single pore, IF—incomplete fusion, V—voids, HQ—high quality, LF—lack of fusion. \* energy input is calculated by simply dividing the laser power by scanning speed, while the heat input should also include the heat transfer efficiency. Other process parameters; powder feeding rate: 8.5 g/min, diameter of the nozzle tip for nitro-

**Laser Power (W)** 

**Scanning Speed (mm/min)** 

gen vapours stream: 5.0 mm, average consumption of liquid nitrogen: 80 ÷ 100 g/min.

*Materials* **2021**, *14*, x FOR PEER REVIEW 6 of 26

4 and 5.

single bead approx. 2.0 ÷ 3.0 mm, and proper fusion to the steel substrate, Table 3, Figures

**Energy Input \* (J/mm)** 

**Remarks Free Cooling/Hybrid** 

**Figure 5.** (**a**) A view of the experimental setup for laser powder deposition of composite layers with forced local cryogenic cooling by the nitrogen vapours stream (from left: cryogenic dewar with automatic control unit, elastic hose for nitrogen delivers, nitrogen stream nozzle and laser head with coaxial nozzles for powder injection); (**b**) a scheme of the laser powder deposition with forced cooling of the substrate by localized nitrogen vapours stream. **Figure 5.** (**a**) A view of the experimental setup for laser powder deposition of composite layers with forced local cryogenic cooling by the nitrogen vapours stream (from left: cryogenic dewar with automatic control unit, elastic hose for nitrogen delivers, nitrogen stream nozzle and laser head with coaxial nozzles for powder injection); (**b**) a scheme of the laser powder deposition with forced cooling of the substrate by localized nitrogen vapours stream. stream of liquid nitrogen. The temperature of the liquid nitrogen in the container was below −196°C, while the temperature of the nitrogen vapours stream measured at the tip of the nozzle by a thermocouple was approx. −160÷164°C. In such a way deep cryogenic treatment (DCT) was provided in the localized area of forced cooling, Table 3, Figure 4b.

**Figure 6.** A view of stringer beads produced by laser deposition of experimental composite WC-Ni powder at free cooling conditions (conventional laser cladding); (**a**) and forced localized cooling by nitrogen vapours stream under cryogenic conditions (hybrid laser deposition process); (**b**), Table 3. **Figure 6.** A view of stringer beads produced by laser deposition of experimental composite WC-Ni powder at free cooling conditions (conventional laser cladding); (**a**) and forced localized cooling by nitrogen vapours stream under cryogenic conditions (hybrid laser deposition process); (**b**), Table 3.

The scanning speed was kept constant at 500 mm/min, while the laser output power was varied from 500 W, 1000 W, 1500 W, and 2000 W. The feeding rate of the composite WC-Ni powder was maintained constant at 8.5 g/min. The powder was fed into the melt pool by three coaxial nozzles with a diameter of 0.8 mm each. The nozzles were attached to head body providing the ability to change the angle of inclination and vertical shift, as can be seen in Figure 4a. The individual streams from the nozzles were focused on the melt pool surface in the region of laser beam interaction, Figure 4. The powder feeding rate was controlled by a rotary disc feeder PFU4 (Durum, Willich, Germany), with the chamber filled in by argon at the pressure 1.5 bar, used also as the carrier gas. The flow The scanning speed was kept constant at 500 mm/min, while the laser output power was varied from 500 W, 1000 W, 1500 W, and 2000 W. The feeding rate of the composite WC-Ni powder was maintained constant at 8.5 g/min. The powder was fed into the melt pool by three coaxial nozzles with a diameter of 0.8 mm each. The nozzles were attached to head body providing the ability to change the angle of inclination and vertical shift, as can be seen in Figure 4a. The individual streams from the nozzles were focused on the melt pool surface in the region of laser beam interaction, Figure 4. The powder feeding rate was controlled by a rotary disc feeder PFU4 (Durum, Willich, Germany), with the chamber filled in by argon at the pressure 1.5 bar, used also as the carrier gas. The flow rate of the carrier argon gas was kept at 8.0 L/min.

rate of the carrier argon gas was kept at 8.0 l/min. The laser beam was set perpendicularly to the surface of the substrate, and it was The laser beam was set perpendicularly to the surface of the substrate, and it was transmitted through a 10.0 mm diameter cylindrical nozzle in argon atmosphere, at the flow

applied configurations of optics. In order to provide wider area of laser beam interaction on the substrate material, the beam was defocused by lifting the laser head and thus focal plane over the top surface of the substrate. In this way, as a result of lifting the focal plane

transmitted through a 10.0 mm diameter cylindrical nozzle in argon atmosphere, at the flow rate 20.0 L/min. The argon flow was used mainly for protection the optics inside the rate 20.0 L/min. The argon flow was used mainly for protection the optics inside the laser head against spatter. While the task of shielding the melt pool was negligible due to long distance (160 mm) between the nozzle and the substrate surface, as a result of defocusing the laser beam spot, Figure 4b. The laser beam spot diameter was 300 µm at the applied configurations of optics. In order to provide wider area of laser beam interaction on the substrate material, the beam was defocused by lifting the laser head and thus focal plane over the top surface of the substrate. In this way, as a result of lifting the focal plane to the distance of 160 mm from the substrate, the diameter of the laser beam was approximately 3.0 mm on the top surface, Figures 4b and 6.

Nitrogen vapours from a pressurized container (cryogenic dewar) were transported via a flexible hose into the jet nozzle. The nozzle tip was designed to be replaceable. Thanks to that different diameter tip nozzles could be applied. In this study a nozzle tip with a diameter of 5.0 mm was used, providing localized cooling of the clad area just behind the melt pool. The jet nozzle was fixed on the laser processing head, and the distance between the nozzle tip and the target surface was maintained constant at 40.0 mm, Figures 4b and 5a. The outflow rate of the nitrogen vapours stream from the nozzle depended on the pressure inside the container, which was maintained constant by an automatic pressure control unit KRIOSAN (KrioSystem, Wrocław, Poland), Figure 5. The average consumption of liquid nitrogen was ranged in 80 ÷ 100 g/min.

#### *2.3. Macro and Microstructure Examinations*

The specimens with test clad layers produced both by conventional laser cladding at free cooling conditions in ambient air, and the clad layers produced at forced cooling under cryogenic conditions were examined by visual inspection first. All the clad layers fulfilled the main criterion of complete penetration of the substrate. Therefore, all the test layers were cut perpendicularly and samples for metallographic study were prepared. Three sliced sections were taken from each test clad layer. One was taken from the middle region, and next two at a distance 15 mm from the beginning and the end of the clad layer, Figure 6. The sliced sections were first mounted in thermosetting phenol resin with graphite filler Electro-WEM (Metalogis, Warsaw, Poland), and next the samples were wet grinded by water papers with grit 120 to 2500 using an automatic grinding/polishing machine Struers Labopol-2 (Struers, Rodovre, Denmark). Next the cross-sections were polished with 1 µm diamond suspension Metkon Diapat-M (Metkon Instruments Inc., Bursa, Turkey). After polishing, the cross-sections were etched by the HNO<sup>3</sup> + 3HCL reagent to disclose the microstructure.

Observations of macrostructure at low magnifications (up to 25×) were carried out by means of a stereoscopic microscope OLYMPUS SZX9 (Olumpus Corporation, Tokyo, Japan), while the microstructure observations were done by the inverted metallographic microscope NIKON Eclipse MA100 (Nikon Corporation, Tokyo, Japan). The microstructure was additionally examined by scanning electron microscopy SEM (Carl Zeiss, Oberkochen, Germany), equipped with the Energy Dispersive Spectrometer EDS (Oxford Instruments, Abingdon, GB, USA) and Phenom Pro-X SEM equipped with EDS and BSD detectors (Thermo Fisher Scientific, Eindhoven, The Netherlands). The phase composition was determined by X-Ray diffraction (Panalitycal, Almelo, The Netherlands) with CuKα source of radiation, with the scanning range of the diffraction angle 2θ from 0 to 140◦ .

#### *2.4. Hardness Measerements*

The hardness was measured on the cross-section of the test clad layers after microstructural studies by Vickers test, at the load 5 N and the dwell time 10 s. The hardness tester WILSON WOLPERT 401 MVD (Wolpert Wilson Instruments, Aachen, Germany) was applied in the study. Measurements were taken along the vertical axis of symmetry, starting from the under surface region of the clad layer (face of the clad). The first measuring point was 0.15 mm under the top surface, while the distance between the subsequent points was

constant at 0.2 mm. In such a way, the hardness distribution from the under surface region, through the clad, fusion zone, heat affected zone, till to the base metal was determined.

#### *2.5. Tribological Test*

The tribological tests of coatings produced by laser deposition of experimental composite powder at free cooling and under cryogenic conditions were conducted by a ball-on-disc tribometer T-01M under room temperature of 22 ◦C, according to the ASTM G99 standard. The relative humidity was about 45% ± 5%. The specimens were prepared in a form of discs with diameter of 45.0 mm balls made of bearing steel (EN 100Cr6, AISI 5210) with a diameter of 10.0 mm were used as the counter face material. The normal load was set as 20 N. The number of revolutions was 1500, while radius of the track was 15 mm. Therefore, the sliding distance was 141.3 m, while the sliding speed was 0.157 m/s. The tangential force of friction and displacement value were continuously measured and recorded during tests using a data acquisition system with PC computer. While the coefficient of friction µ was calculated by dividing the value of tangential force of friction by the value of normal load used:

$$
\mu = \mathbf{T} / \mathbf{F}\_{\mathbf{n}\nu} \tag{1}
$$

where:

Fn—normal load (30 N), T—tangential force of friction.

#### **3. Results and Discussion**

#### *3.1. Macrostructure and Single Bead Geometry*

At the first stage of the study the cross-sections of the clad layers produced as single stringer beads by conventional laser cladding at free cooling conditions were compared with the single stringer beads produced by the novel technique considered as a hybrid laser deposition under cryogenic conditions, Figure 7 Since the clad layers were produced at constant scanning speed, the influence of energy input on the bead geometry can be also considered. The energy input is defined as follows:

$$\mathbf{Ev} = \mathbf{P}/\mathbf{v} \text{ (J/mm)},\tag{2}$$

where P is the output power of laser beam (W), and v is the scanning speed (mm/s).

It is worth to note, when the heat transfer efficiency into the material is known and considered in the calculation, this parameter is called heat input. In laser processing the heat transfer efficiency is related with the absorption of laser energy, and under certain conditions it can take the value 1. Therefore, for simplicity the term "heat input" will be used in the remainder of the manuscript text. As seen in Figure 7, clear interface lines between the clads and the substrate can be distinguished, as well as the heat affected zone regions. A rough comparison of the geometry of clad layers produced at the same heat input but different cooling conditions, indicates that the forced localized cooling by nitrogen vapours stream leads to a clear decrease of penetration depth and thus dilution.

The value of dilution "*D*" was calculated by the following formula:

$$D = \frac{A\_{FZ}}{A\_{FZ} + A\_{CL}} \cdot 100\%\_{\text{\textdegree}} \tag{3}$$

where: *AFZ* is the cross-section area of the fusion zone, and the *ACL* is the cross-section area of the clad layer (Figure 7).

The increase in laser output power, and therefore increase in heat input, leads to an increase in the curvature of the interface line, penetration depth, height and width of the clad, Figures 7 and 8.

on-disc tribometer T-01M under room temperature of 22 °C, according to the ASTM G99 standard. The relative humidity was about 45% ± 5%. The specimens were prepared in a form of discs with diameter of 45.0 mm balls made of bearing steel (EN 100Cr6, AISI 5210) with a diameter of 10.0 mm were used as the counter face material. The normal load was set as 20 N. The number of revolutions was 1500, while radius of the track was 15 mm. Therefore, the sliding distance was 141.3 m, while the sliding speed was 0.157 m/s. The tangential force of friction and displacement value were continuously measured and recorded during tests using a data acquisition system with PC computer. While the coefficient of friction µ was calculated by dividing the value of tangential force of friction by

At the first stage of the study the cross-sections of the clad layers produced as single stringer beads by conventional laser cladding at free cooling conditions were compared with the single stringer beads produced by the novel technique considered as a hybrid laser deposition under cryogenic conditions, Figure 7 Since the clad layers were produced at constant scanning speed, the influence of energy input on the bead geometry can be

It is worth to note, when the heat transfer efficiency into the material is known and considered in the calculation, this parameter is called heat input. In laser processing the heat transfer efficiency is related with the absorption of laser energy, and under certain conditions it can take the value 1. Therefore, for simplicity the term "heat input" will be used in the remainder of the manuscript text. As seen in Figure 7, clear interface lines between the clads and the substrate can be distinguished, as well as the heat affected zone

input but different cooling conditions, indicates that the forced localized cooling by nitrogen vapours stream leads to a clear decrease of penetration depth and thus dilution.

*Materials* **2021**, *14*, x FOR PEER REVIEW 9 of 26

where P is the output power of laser beam (W), and v is the scanning speed (mm/s).

µ=T/Fn, (1)

Ev = P/v (J/mm), (2)

the value of normal load used:

**3. Results and Discussion** 

*3.1. Macrostructure and Single Bead Geometry* 

Fn—normal load (30 N), T—tangential force of friction.

also considered. The energy input is defined as follows:

where:

**Figure 7.** Macrostructure and the single bead geometry of the clad layers produced by laser cladding of experimental composite WC-Ni powder at constant scanning speed 500 mm/min, Table 3: (**a**) output power 500 W at free cooling; (**b**) output power 500 W at forced localized cooling; (**c**) output power 2000 W at free cooling; (**d**) 2000 W at forced localized cooling. **Figure 7.** Macrostructure and the single bead geometry of the clad layers produced by laser cladding of experimental composite WC-Ni powder at constant scanning speed 500 mm/min, Table 3: (**a**) output power 500 W at free cooling; (**b**) output power 500 W at forced localized cooling; (**c**) output power 2000 W at free cooling; (**d**) 2000 W at forced localized cooling. of the clad layer (Figure 7). The increase in laser output power, and therefore increase in heat input, leads to an increase in the curvature of the interface line, penetration depth, height and width of the clad, Figures 7 and 8.

**Figure 8.** Influence of the process parameters and cooling conditions on the geometry and dimensions of the stringer beads produced by laser deposition of WC-Ni powder (Table 3): (**a**) height of the clad (reinforcement); (**b**) penetration depth into the substrate; (**c**) width of the single clad layer; and (**d**) calculated rate of dilution. **Figure 8.** Influence of the process parameters and cooling conditions on the geometry and dimensions of the stringer beads produced by laser deposition of WC-Ni powder (Table 3): (**a**) height of the clad (reinforcement); (**b**) penetration depth into the substrate; (**c**) width of the single clad layer; and (**d**) calculated rate of dilution.

(c) (d) **Figure 8.** Influence of the process parameters and cooling conditions on the geometry and dimensions of the stringer beads produced by laser deposition of WC-Ni powder (Table 3): (**a**) height of the clad (reinforcement); (**b**) penetration depth into

the substrate; (**c**) width of the single clad layer; and (**d**) calculated rate of dilution.

Based on the visual inspection and analysis of cross-sections morphology, no cracks neither tendency for cracking was observed both in the clad layers produced at free cooling and under cryogenic conditions. It is related with the relative high ductility of the nickel base matrix. However, single pores were found on cross-sections of the clads produced at free cooling in the range of heat input 120 ÷ 240 J/mm. The single pore with the highest diameter of 180 µm was found on the cross-section of the clad LC3. In turn, the tendency for porosity was found to be less in the case of clads produced under cryogenic conditions. Just a single pore was found just on the cross-section of the clad produced at the highest heat input of 240 J/mm but its diameter was lower, approx. 140 µm. Thorough analysis of the geometrical parameters on cross-sections of the clads showed that the localized forced cooling of the cladding region by a stream of nitrogen vapours has neglectable influence on the height of the clads. As can be seen in Figure 8a, the height of the clads produced under different conditions of cooling (free cooling and localized forced cryogenic cooling) is comparable. In turn, the average width of the clads produced under cryogenic conditions is approximately 5 ÷ 15% lower if compared to the clads produced at free cooling (conventional laser cladding), Figure 8c. A greater difference can be noticed in the case of penetration depth and dilution, Figure 8b,d. The average penetration depth of the clads produced at the same heat input but localized forced cooling (cryogenic conditions) is 24% to 38% lower than in the case of clads produced at free cooling conditions. While the average dilution of the clads produced at localized forced cooling is 20 ÷ 25% lower if compared to the clad produced by conventional laser cladding at free cooling.

The obtained results indicate that the localized cooling of the substrate in the region of cladding can effectively reduce the penetration depth by dissipating heat. However, the applied technique of localized cooling does not significantly reduce the height and width of the clads, thus also the area of the clads. Simultaneously, the applied technique of localized cooling has a beneficial effect on reducing the penetration depth and dilution, as can be seen in Figure 8b,d. This phenomenon is caused by the localized cooling just a narrow region along the axis of the stringer bead which leads to reducing the maximum temperature in the middle region of laser beam interaction and reduces the temperature gradient.

#### *3.2. Microstructure*

The comparative microstructures of the clad layers produced at maximum heat input of 240 J/mm (output power 2000 W, scanning speed 500 mm/min) and different cooling conditions (LC4 and HC4) are presented on optical micrographs in Figure 9, while the microstructures of the clad layers produced at minimum heat input of 60 J/mm and different cooling conditions (LC1 and HC1) are presented on optical micrographs in Figure 10. Comparing the Figure 9a,b, it can be seen, that the number and size of the carbides are different on the cross sections of the clads produced at the same heat input but different cooling conditions. Only single massive carbides can be found on the cross section of the clad layer produced at maximum heat input and free cooling conditions. In the case of the clad layer produced at the same heat input but at forced cooling by the stream of nitrogen vapours, the number of carbides is clearly higher. Closer view of the microstructure of the clad produced at maximum heat input and free cooling conditions (LC4) revealed that the regions between massive carbides are rich mainly in needle-like precipitations with a length up to approx. 40 µm, Figure 9c. For comparison, the microstructure in the regions between massive carbides of the clad produced at the maximum heat input but at localized forced cooling is different, Figure 9c. As can be seen, the number and size of precipitation is much smaller. The longest needle-like precipitations have a length up to approx. 15 µm, Figure 9d. The reason of this phenomenon is different cooling thus different thermal conditions during deposition and solidifying of the two comparative clad layers.

**Figure 9.** Microstructure of the comparative stringer beads LC4 and HC4 produced at the same maximum heat input of 240 J/mm (scanning speed 500 mm/min, laser output power 2000 W) but different cooling conditions (Table 3): (**a**,**c**) free cooling in ambient air—a view of the entire clad layer and the middle region of the clad respectively; and (**b**,**d**) forced localized cooling by nitrogen vapours stream under cryogenic conditions ("hybrid laser deposition")—a view of the entire clad layer and the middle region of the clad respectively. **Figure 9.** Microstructure of the comparative stringer beads LC4 and HC4 produced at the same maximum heat input of 240 J/mm (scanning speed 500 mm/min, laser output power 2000 W) but different cooling conditions (Table 3): (**a**,**c**) free cooling in ambient air—a view of the entire clad layer and the middle region of the clad respectively; and (**b**,**d**) forced localized cooling by nitrogen vapours stream under cryogenic conditions ("hybrid laser deposition")—a view of the entire clad layer and the middle region of the clad respectively.

> The clad layer LC4 was produced by conventional laser cladding at free cooling in ambient and at relatively high heat input 240 J/mm, Table 3. Under such conditions, the massive carbides exposed to high temperature for a relatively long time, decompose and dissolve in the melt pool. It is worth noting that in addition to being exposed to the thermal influence of the melt pool, carbides are also exposed to the direct interaction of the laser beam radiation. The clad layer LC4 was produced by conventional laser cladding at free cooling in ambient and at relatively high heat input 240 J/mm, Table 3. Under such conditions, the massive carbides exposed to high temperature for a relatively long time, decompose and dissolve in the melt pool. It is worth noting that in addition to being exposed to the thermal influence of the melt pool, carbides are also exposed to the direct interaction of the laser beam radiation.

> The carbide particles dissolved in the melt pool first enrich the liquid solution with carbon and tungsten and then recrystalize as secondary carbides in the form of dendrites, needle-like or block-like precipitations. So, the optical micrographs of the clad layer LC4 are typical for the clad produced at excessive energy and heat input, characterized by unfavourable microstructure with low share of massive tungsten carbides and high amount of secondary precipitations, Figure 9a,c. On the other hand, the clad layer HC4 produced at the same processing parameters and heat input 240 J/mm but additionally with the localized cooling showed clearly lower tendency for decomposition the massive carbides due to dissipation of the excessive heat by the nitrogen vapours stream under cryogenic conditions. Therefore, the amount and size of precipitations in the regions between massive carbides is lower. The massive carbides show "feather" structure, typical for eutectic W2C/WC carbides, Figure 9d. The carbide particles dissolved in the melt pool first enrich the liquid solution with carbon and tungsten and then recrystalize as secondary carbides in the form of dendrites, needle-like or block-like precipitations. So, the optical micrographs of the clad layer LC4 are typical for the clad produced at excessive energy and heat input, characterized by unfavourable microstructure with low share of massive tungsten carbides and high amount of secondary precipitations, Figure 9a,c. On the other hand, the clad layer HC4 produced at the same processing parameters and heat input 240 J/mm but additionally with the localized cooling showed clearly lower tendency for decomposition the massive carbides due to dissipation of the excessive heat by the nitrogen vapours stream under cryogenic conditions. Therefore, the amount and size of precipitations in the regions between massive carbides is lower. The massive carbides show "feather" structure, typical for eutectic W2C/WC carbides, Figure 9d.

> Beside comparison of microstructure, the optical micrographs revealed a cluster of pores on the cross section of the clad LC4 produced by conventional laser cladding. The

20 ÷ 30 µm.

produced under forced cooling conditions, Figure 10d.

**Figure 10.** Microstructure of the comparative stringer beads LC1 and HC1 produced at the same minimum heat input of 60 J/mm (scanning speed 500 mm/min, laser output power 500 W) but different cooling conditions (Table 3): (**a**,**c**) free cooling in ambient air—a view of the entire clad layer and the middle region of the clad respectively; and (**b**,**d**) forced localized cooling by nitrogen vapours stream under cryogenic conditions ("hybrid laser deposition")—a view of the entire clad layer and the middle region of the clad respectively. **Figure 10.** Microstructure of the comparative stringer beads LC1 and HC1 produced at the same minimum heat input of 60 J/mm (scanning speed 500 mm/min, laser output power 500 W) but different cooling conditions (Table 3): (**a**,**c**) free cooling in ambient air—a view of the entire clad layer and the middle region of the clad respectively; and (**b**,**d**) forced localized cooling by nitrogen vapours stream under cryogenic conditions ("hybrid laser deposition")—a view of the entire clad layer and the middle region of the clad respectively.

Beside comparison of microstructure, the optical micrographs revealed a cluster of pores on the cross section of the clad LC4 produced by conventional laser cladding. The pores are placed in the upper part of the clad layer and the average diameter is approx. 20 ÷ 30 µm.

pores are placed in the upper part of the clad layer and the average diameter is approx.

In turn, the optical micrographs of the comparative clad layers produced at minimum heat input of 60 J/mm and different cooling conditions (LC1 and HC1) also show differences in microstructure, Figure 10. First, if compared to the clad layers produced at the maximum heat input, these clad layers have more massive carbides on cross sections and their share in relation to the matrix is significantly greater, Figure 10. Additionally, the carbides are evenly distributed on the cross sections. However, the number and share of massive carbides is clearly greater in the case of the clad layer produced under forced cooling by the localized stream of nitrogen vapours, Figure 10a,b. Observations of the optical micrographs confirmed the tendency for porosity in the case of clad layers produced by conventional laser cladding at free cooling conditions, even at the minimum heat input of 60 J/mm, Figure 10a. The cluster of pores can be observed in the middle part of the clad layer LC1, while the average diameter of pores ranges between 20 to 80 µm. In turn, just two single pores with a diameter approx. 15 µm were observed on cross section of the clad

In turn, the optical micrographs of the comparative clad layers produced at minimum heat input of 60 J/mm and different cooling conditions (LC1 and HC1) also show differences in microstructure, Figure 10. First, if compared to the clad layers produced at the maximum heat input, these clad layers have more massive carbides on cross sections and their share in relation to the matrix is significantly greater, Figure 10. Additionally, the carbides are evenly distributed on the cross sections. However, the number and share of massive carbides is clearly greater in the case of the clad layer produced under forced cooling by the localized stream of nitrogen vapours, Figure 10a,b. Observations of the optical micrographs confirmed the tendency for porosity in the case of clad layers produced by conventional laser cladding at free cooling conditions, even at the minimum heat input of 60 J/mm, Figure 10a. The cluster of pores can be observed in the middle part of the clad layer LC1, while the average diameter of pores ranges between 20 to 80 µm. In turn, just two single pores with a diameter approx. 15 µm were observed on cross section of the clad produced under forced cooling conditions, Figure 10d.

In contrast to the clad layer produced at maximum heat input and free cooling, the microstructure of the clad LC1 in the regions between the massive carbides is characterized by a low number of individual precipitations not exceeding 10 µm in length, as can be clearly seen in Figure 10c. There are two explanations for this phenomenon. The first is related with lower amount of decomposed and diluted carbides due to four times lower laser output power (500 W) and thus the heat input. Therefore, the enrichment of the liquid by the W and C elements was also limited and insufficient for precipitation of carbides. The second reason is related with lower volume of the melt pool thus shorter time of material being at liquid state and shorter time for precipitation.

Close observation of the massive carbide's morphology on the optical micrographs in Figure 10d, showed that all these carbides are covered with a thin layer of block-like particles. Similar observations were reported by Zhang P. et al. [12] for spherical and irregularly shaped carbides in nickel-based matrix. Jones M. and Waag U. provided a comprehensive study on the effect of WC particle type on Ni-based composite coatings. They explained this phenomenon by partial melting of the WC particles during laser cladding. Next the C and W elements combine with Cr, Fe, and Ni to form the complex carbides M23C6, M7C3, and M6C, which distribute in the matrix to form needle-like or block-like. However, a detailed analysis of this boundary and an attempt to explain this phenomenon is provided in the further part of the discussion.

To quantify the relationship between of laser output power of deposition process (proportional to heat input at constant scanning speed) and the cooling conditions and the morphology of the clad layer, the share of carbides relative to the matrix was determined on the cross sections, and the carbides size distribution for different clad layers was determined. The calculations were made by digitalising the micrographs images and using the NIS-Elements software (Nikon Corporation, Tokyo, Japan) for determining the area share of carbides, and the results are presented in Figure 11.

As can be seen, the calculated percentage share of carbides on cross section of the clad layer produced at maximum heat input of 240 J/mm and free cooling conditions is below 7%, while the share of carbides in the case of clad layer produced at the same heat input but with additional forced cooling is significantly higher, approx. over 22%, Figure 11a.

The relationship indicates generally the lower heat input the higher the share of carbides in the matrix. However, as can be seen, the effect of the forced localized cooling is advantageous, because it provides significantly higher share of carbides for every clad layer produced in the range of investigated parameters. In turn, the carbides size distribution determined for the comparative clad layers produced at minimum heat input of 60 J/mm show that the dominant size for the clad layer LC1 produce under free cooling conditions is approx. 150 ÷ 200 µm. While in the case of the clad layer HC1 produce under forced cooling conditions the carbides distribution is different and the share of carbides sized in the range 50 up to 150 µm is even and proportional, Figure 11b. The relationship clearly indicates that the cooling conditions at the same heat input affect significantly the thermal conditions and the tendency to dissolve carbides and thus reduce their size is lower when forced cooling is applied.

In turn the size distribution of carbides determined on the surface of the specimen indicates that the proportion of fine carbides is significant, higher than it would appear from the fraction share for the powder. The first reason for this phenomenon is the presence of secondary carbides precipitated in the liquid because of crystallization. However, these are carbides mainly in dendritic form. The share of such dendritic carbides is greater in the range of higher heat inputs (laser power), due to the conditions conducive to the dissolve of carbides in the liquid solution because of diffusion, decomposition due to direct heating with a laser beam, and then precipitation from the liquid solution. However, the presence of a large amount of very small carbide particles is also evident. The explanation for the presence of such many fine carbides is the primary form and morphology of WC-type carbides, which are in fact the eutectic of WC + W2C manufactured as a result of melting

*Materials* **2021**, *14*, x FOR PEER REVIEW 13 of 26

of material being at liquid state and shorter time for precipitation.

is provided in the further part of the discussion.

and crushing. As can be seen from Figures 10d, 12–14 and 15a,b these massive eutectic carbides are coated by a layer of stoichiometric WC with high thermal stability. area share of carbides, and the results are presented in Figure 11.

In contrast to the clad layer produced at maximum heat input and free cooling, the microstructure of the clad LC1 in the regions between the massive carbides is characterized by a low number of individual precipitations not exceeding 10 µm in length, as can be clearly seen in Figure 10c. There are two explanations for this phenomenon. The first is related with lower amount of decomposed and diluted carbides due to four times lower laser output power (500 W) and thus the heat input. Therefore, the enrichment of the liquid by the W and C elements was also limited and insufficient for precipitation of carbides. The second reason is related with lower volume of the melt pool thus shorter time

Close observation of the massive carbide's morphology on the optical micrographs in Figure 10d, showed that all these carbides are covered with a thin layer of block-like particles. Similar observations were reported by Zhang P. et al. [12] for spherical and irregularly shaped carbides in nickel-based matrix. Jones M. and Waag U. provided a comprehensive study on the effect of WC particle type on Ni-based composite coatings. They explained this phenomenon by partial melting of the WC particles during laser cladding. Next the C and W elements combine with Cr, Fe, and Ni to form the complex carbides M23C6, M7C3, and M6C, which distribute in the matrix to form needle-like or block-like. However, a detailed analysis of this boundary and an attempt to explain this phenomenon

To quantify the relationship between of laser output power of deposition process (proportional to heat input at constant scanning speed) and the cooling conditions and the morphology of the clad layer, the share of carbides relative to the matrix was determined on the cross sections, and the carbides size distribution for different clad layers was determined. The calculations were made by digitalising the micrographs images and using the NIS-Elements software (Nikon Corporation, Tokyo, Japan) for determining the

indicates that the proportion of fine carbides is significant, higher than it would appear from the fraction share for the powder. The first reason for this phenomenon is the pres-

in the range of higher heat inputs (laser power), due to the conditions conducive to the dissolve of carbides in the liquid solution because of diffusion, decomposition due to direct heating with a laser beam, and then precipitation from the liquid solution. However, the presence of a large amount of very small carbide particles is also evident. The explanation for the presence of such many fine carbides is the primary form and morphology of WC-type carbides, which are in fact the eutectic of WC + W2C manufactured as a result of melting and crushing. As can be seen from Figures 10d, 12–15a,b these massive eutectic

However, in the conditions of laser cladding the particles are exposed to direct action of the laser beam, and thus separation or even decomposition. If the carbide gets directly to the melt pool with a temperature below the carbides melting point, its decomposition will be slight, mainly due to the mutual diffusion of carbide and the liquid. When the carbide is exposed to direct laser beam radiation with high power density, it can directly

carbides are coated by a layer of stoichiometric WC with high thermal stability.

In turn the size distribution of carbides determined on the surface of the specimen **Figure 12.** SEM micrographs of the comparative stringer beads LC1 and HC1 produced at the same minimum heat input of 60 J/mm (scanning speed 500 mm/min, laser output power 500 W) but different cooling conditions (Table 3): (**a**) free **Figure 12.** SEM micrographs of the comparative stringer beads LC1 and HC1 produced at the same minimum heat input

clad layer and the middle region of the clad respectively.

cooling in ambient air—a view of the entire clad layer and the middle region of the clad respectively; and (**b**) forced

of 60 J/mm (scanning speed 500 mm/min, laser output power 500 W) but different cooling conditions (Table 3): (**a**) free cooling in ambient air—a view of the entire clad layer and the middle region of the clad respectively; and (**b**) forced localized cooling by nitrogen vapours stream under cryogenic conditions ("hybrid laser deposition")—a view of the entire clad layer and the middle region of the clad respectively. *Materials* **2021**, *14*, x FOR PEER REVIEW 16 of 26

**Figure 13.** SEM microstructure of the stringer bead HC1 produced at the minimum heat input of 60 J/mm (scanning speed 500 mm/min, laser output power 500 W) and forced localized cooling by nitrogen vapours stream under cryogenic conditions ("hybrid laser deposition", Table 3): (**a**) general view of massive carbides and matrix; (**b**) a view of carbides boundaries; (**c**) partially recristalized needle-like carbides; (**d**) block-like particles in the dendritic matrix. **Figure 13.** SEM microstructure of the stringer bead HC1 produced at the minimum heat input of 60 J/mm (scanning speed 500 mm/min, laser output power 500 W) and forced localized cooling by nitrogen vapours stream under cryogenic conditions ("hybrid laser deposition", Table 3): (**a**) general view of massive carbides and matrix; (**b**) a view of carbides boundaries; (**c**) partially recristalized needle-like carbides; (**d**) block-like particles in the dendritic matrix.

Another evidence that the light grey particles on the boundaries of massive eutectic carbides were not recrystalized are the very fine crystals growing epitaxially on the tips **Table 4.** Summary of EDS analysis on the cross section of the stringer bead HC1 produced at the minimum heat input of 60 J/mm and forced cryogenic cooling conditions (hybrid laser deposition process), Figure 14.


metastable almost up to melting point. However, it should be noted that the conditions of phases transformation during laser deposition are far from the conditions of thermody-

In the case of chemical composition analysis with an EDS spectrometer, the quantitative content of elements such as nitrogen and carbon are usually not given due to the


**Table 4.** *Cont.*

ensure greater clarity of the results.

*Materials* **2021**, *14*, x FOR PEER REVIEW 17 of 26

high uncertainty and measurement error. However, in the case of the presented results of the study on chemical composition, a satisfactory and sufficient convergence of the results of carbon and nitrogen content was obtained, allowing for the identification and confirmation of the presence of the expected constituents. For this reason, it was decided to include the results of the quantitative measurements of carbon and nitrogen content to

The EDS analysis conducted on the cross-section of the clad layer HC1 produced at the lowest heat input and forced cooling, confirmed that the composition of the massive carbides is typical for eutectic W2C/WC, Figure 14, Table 4. In turn, the composition of the

**Figure 14.** SEM microstructure of the selected representative regions between the massive tungsten carbides on cross section of the stringer bead HC1 produced at the minimum heat input of 60 J/mm (scanning speed 500 mm/min, laser output power 500 W) and forced localized cooling by nitrogen vapours stream under cryogenic conditions ("hybrid laser deposition", Table 3): (**a**,**b**) a view of the region between massive carbides; results of EDS for individual points and line scan provided in Table 4. **Figure 14.** SEM microstructure of the selected representative regions between the massive tungsten carbides on cross section of the stringer bead HC1 produced at the minimum heat input of 60 J/mm (scanning speed 500 mm/min, laser output power 500 W) and forced localized cooling by nitrogen vapours stream under cryogenic conditions ("hybrid laser deposition", Table 3): (**a**,**b**) a view of the region between massive carbides; results of EDS for individual points and line scan provided in Table 4.

**Table 4.** Summary of EDS analysis on the cross section of the stringer bead HC1 produced at the minimum heat input of 60 J/mm and forced cryogenic cooling conditions (hybrid laser deposition process), Figure 14. **Table 5.** Summary of EDS analysis on the cross section of the stringer bead LC1 produced at the minimum heat input of 60 J/mm and free cooling conditions (conventional laser cladding), Figure 15.


bides of clad layer produced at forced cooling, Figures 14 and 15.

a\_6 56.4/94.3 1.8/0.9 - - 42.0/4.7 - - WC b\_1 58.1/94.6 2.0/1.0 - - 39.9/4.4 - - WC b\_2 60.9/91.9 10.0/5.0 - - 29.1/3.1 - - M2C b\_3 55.7/90.2 11.5/6.2 - - 32.8/3.6 - - M2C b\_4 59.7/92.7 7.1/3.7 - - 33.2/3.7 - - M2C b\_5 21.1/54.0 49.8/40.7 - - 21.9/3.7 7.2/1.6 - Primary γ-Ni b\_6 21.0/54.5 48.3/39.9 - - 23.2/3.9 7.5/1.7 - Primary γ-Ni b\_7 6.0/25.7 37.8/52.3 4.5/2.9 0.9/1.1 17.4/5.0 27.7/10.6 3.0/2.36 Ni/Ni3B eutectic,

b\_8 9.9/33.3 54.4/58.3 - - 26.8/5.9 - 8.8/2.6 Ni/Ni3B eutectic,

Although the quantitative results of the measurement of nitrogen content should be considered with a great uncertainty, it is evident that the nitrogen was detected in the clad layer produced under cryogenic conditions, mainly in regions of matrix dendrites and interdendritic regions, Figure 14 and Table 4. It is most likely the result of nitrogen vapours stream application directly into the deposition region, and thus the presence of high amount of gaseous nitrogen in this region. In any case, the test results show that there is

On the other hand, the morphology of the clad layer LC1 produced at minimum heat input but at free cooling conditions observed on the SEM micrographs is completely different that the morphology of the clad layer produced under forced cooling conditions, Figure 15. Beside significantly lower share of massive eutectic carbides, the morphology of precipitations in the regions between the massive eutectic carbides is also different. Closer observation of the massive carbides revealed another difference in the morphology. These carbides observed on the cross-section of the clad layer LC1 produced at free cooling have been dissolved much less in the boundary regions if compared with the car-

partial nitrogen absorption under the investigated conditions of laser deposition.

b\_9 64.4/93.6 7.7/3.7 - - 27.9/2.7 - - M7C3 b\_10 32.1/69.7 36.8/25.4 - - 20.1/2.8 11.0/2.1 - Line scan

M2C

M2C

**Figure 15.** SEM microstructure of the stringer bead LC1 produced at the minimum heat input of 60 J/mm (scanning speed 500 mm/min, laser output power 500 W) and free cooling in ambient air (Table 3): (**a**,**b**) general view of massive carbides and matrix; (**c**) a view of carbides boundaries; (**d**) a view of precipitations in the matrix. Results of EDS for individual points and line scan provided in Table 5. **Figure 15.** SEM microstructure of the stringer bead LC1 produced at the minimum heat input of 60 J/mm (scanning speed 500 mm/min, laser output power 500 W) and free cooling in ambient air (Table 3): (**a**,**b**) general view of massive carbides and matrix; (**c**) a view of carbides boundaries; (**d**) a view of precipitations in the matrix. Results of EDS for individual points and line scan provided in Table 5.

Just little traces of local dissolution and subsequent resolidification can be found on the carbides' boundary. However, it should be noted that these are just few massive eutectic carbides that have not completely dissolved. The explanation is that the liquid solution has reached the solubility limit of carbon and tungsten, at least locally, which inhibited the further dissolution of remaining massive carbides in the melt pool. The resolidified regions of boundaries show fine light grey dendritic precipitation smaller than 10 µm, Figure 15b,d. In turn, the massive carbides are also composed of needle-like particles and However, in the conditions of laser cladding the particles are exposed to direct action of the laser beam, and thus separation or even decomposition. If the carbide gets directly to the melt pool with a temperature below the carbides melting point, its decomposition will be slight, mainly due to the mutual diffusion of carbide and the liquid. When the carbide is exposed to direct laser beam radiation with high power density, it can directly lead to its partial or complete decomposition. In such a case, the carbide decomposes into fine carbides of needle-like or block-like form, as can be seen in Figures 9b and 14a,b.

trapezoidal blocks, however, highly integrated and embedded in the matrix-like phase, as can be seen in Figure 15b. It can be supposed that the matrix-like phase is the sub stoichiometric tungsten carbide W2C (β phase) which is less stable than the stoichiometric monocarbide WC (δ phase). In Figure 15 there are no traces of fine needle-like particles and trapezoidal blocks typical for the clad layer HC1 produced at forced cooling under cryogenic conditions. This is another evidence that these particles do not nuclei nor solidify in the liquid melt pool under the investigated laser cladding conditions. The particles rather come from the decomposition and dissolution of massive eutectic W2C/WC carbides. Since the particles are small and have low heat capacity, thus they heat up intensively and next completely dissolve in the liquid malt pool. However, it must be emphasised that the observations and findings are valid just for the specific conditions of the experiment, determined technological conditions such as the laser beam characteristic (energy density, The representative and comparative SEM micrographs were presented for the clad layers produced at the minimum heat input and different cooling conditions in Figure 12. The SEM micrograph of the clad layer produced at free cooling in Figure 12a show fine dendritic precipitations mainly near the massive carbides. In turn, on the SEM micrograph of the clad layer produced under forced cooling conditions in Figure 12b the regions between the massive carbides are densely filled mainly with fine block-like particles. The SEM micrographs with higher magnification of the clad layer HC1 produced at forced cooling revealed fine dendritic and needle-like particles, as well as trapezoidal blocks evenly distributed in dendritic matrix, Figure 13. While the massive eutectic W2C/WC carbides with feather structure are composed of needle-like and trapezoidal blocks of carbides which are light grey and white, Figure 13a–c. The boundaries of the massive carbides are partially dissolved, as can be seen in Figure 13b,c. It is also characteristic

intensity profile, wavelength, etc.), processing parameters, the way of powder delivery,

presence of M2C type carbides as dendritic precipitations, Figure 15 and Table 5. In turn, the densely packed light grey precipitations between the massive carbides show compo-

sition typical for M3C carbides in eutectic, Figure 15, Table 5.

thermal conditions.

that all needles and trapezoidal blocks on the boundary are light grey. Between the light grey particles at the massive eutectic carbide's boundary the grey dendrites and dark grey interdendritic regions of matrix can be clearly seen, Figure 13d. It can be also seen that some of the light grey particles, especially long needles (length up to 40 µm), go deep into the massive eutectic carbides, Figure 13b,c. The observed and described morphology of the light grey particles (needle-like and trapezoidal blocks) on the boundary of massive eutectic carbides indicates that the particles were not formed because of recrystallization or precipitation from the liquid phase of the melt pool, but the particles are the integrated part of the massive eutectic carbides, Figure 13.

Another evidence that the light grey particles on the boundaries of massive eutectic carbides were not recrystalized are the very fine crystals growing epitaxially on the tips of some of needles and tops of trapezoidal blocks, as shown in Figure 13c,d. Since the needles' tips have been exposed to longest exposure to the liquid and higher temperature, therefore, these regions were dissolved and recrystalized. In turn, smaller separate particles (both needles and blocks) almost completely recrystalized into dendrites, as shown in Figure 13b,d. The above observations indicate that just more thermodynamically stable phases (stoichiometric carbides) remained at the boundary of massive eutectic carbides. Since the massive eutectic carbides (eutectic composition lies at approx. 38 at.% C on W-C phase diagram) are originally composed of approx. 80% of W2C (β phase) and 20% of stoichiometric WC (δ phase), it can be assumed that the remaining phases on boundary are WC carbides. According to a W-C phase diagram, the melting point of tungsten monocarbide WC is at least 50÷60 degree higher if compared to melting point of W2C [32,33]. According to some literature data, the difference is even greater. The melting point for β-W2C is often reported at 2730 ◦C, while for δ-WC at 2870 ◦C. Moreover, the δ-WC phase is metastable almost up to melting point. However, it should be noted that the conditions of phases transformation during laser deposition are far from the conditions of thermodynamic equilibrium, due to high rates of heating, cooling and high maximum temperature and also high temperature gradients.

In the case of chemical composition analysis with an EDS spectrometer, the quantitative content of elements such as nitrogen and carbon are usually not given due to the high uncertainty and measurement error. However, in the case of the presented results of the study on chemical composition, a satisfactory and sufficient convergence of the results of carbon and nitrogen content was obtained, allowing for the identification and confirmation of the presence of the expected constituents. For this reason, it was decided to include the results of the quantitative measurements of carbon and nitrogen content to ensure greater clarity of the results.

The EDS analysis conducted on the cross-section of the clad layer HC1 produced at the lowest heat input and forced cooling, confirmed that the composition of the massive carbides is typical for eutectic W2C/WC, Figure 14, Table 4. In turn, the composition of the trapezoidal block indicates presence of stoichiometric monocrystalline tungsten carbides WC, Figure 14, Table 4. The determined composition of grey dendrites indicates for the matrix of primary γ-Ni, while the dark grey interdendritic regions show composition typical for Ni/Ni3B eutectic, and possible share of M2C type carbides, Figure 14, Table 4. The EDS analysis conducted for the needle-like constituents indicated composition typical for complex carbide M7C<sup>3</sup> type, Figure 14, Table 4. It should be also noted that the content of iron Fe detected in the clad layer LC1 is relatively low. The Fe was confirmed just in the interdendritic regions, Figure 14, Table 4. It indicates that thanks low dilution of the clad by the substrate material, the composition of the clad is close to the applied powder.

Although the quantitative results of the measurement of nitrogen content should be considered with a great uncertainty, it is evident that the nitrogen was detected in the clad layer produced under cryogenic conditions, mainly in regions of matrix dendrites and interdendritic regions, Figure 14 and Table 4. It is most likely the result of nitrogen vapours stream application directly into the deposition region, and thus the presence of high amount of gaseous nitrogen in this region. In any case, the test results show that there is partial nitrogen absorption under the investigated conditions of laser deposition.

On the other hand, the morphology of the clad layer LC1 produced at minimum heat input but at free cooling conditions observed on the SEM micrographs is completely different that the morphology of the clad layer produced under forced cooling conditions, Figure 15. Beside significantly lower share of massive eutectic carbides, the morphology of precipitations in the regions between the massive eutectic carbides is also different. Closer observation of the massive carbides revealed another difference in the morphology. These carbides observed on the cross-section of the clad layer LC1 produced at free cooling have been dissolved much less in the boundary regions if compared with the carbides of clad layer produced at forced cooling, Figures 14 and 15.

Just little traces of local dissolution and subsequent resolidification can be found on the carbides' boundary. However, it should be noted that these are just few massive eutectic carbides that have not completely dissolved. The explanation is that the liquid solution has reached the solubility limit of carbon and tungsten, at least locally, which inhibited the further dissolution of remaining massive carbides in the melt pool. The resolidified regions of boundaries show fine light grey dendritic precipitation smaller than 10 µm, Figure 15b,d. In turn, the massive carbides are also composed of needle-like particles and trapezoidal blocks, however, highly integrated and embedded in the matrix-like phase, as can be seen in Figure 15b. It can be supposed that the matrix-like phase is the sub stoichiometric tungsten carbide W2C (β phase) which is less stable than the stoichiometric monocarbide WC (δ phase). In Figure 15 there are no traces of fine needle-like particles and trapezoidal blocks typical for the clad layer HC1 produced at forced cooling under cryogenic conditions. This is another evidence that these particles do not nuclei nor solidify in the liquid melt pool under the investigated laser cladding conditions. The particles rather come from the decomposition and dissolution of massive eutectic W2C/WC carbides. Since the particles are small and have low heat capacity, thus they heat up intensively and next completely dissolve in the liquid malt pool. However, it must be emphasised that the observations and findings are valid just for the specific conditions of the experiment, determined technological conditions such as the laser beam characteristic (energy density, intensity profile, wavelength, etc.), processing parameters, the way of powder delivery, thermal conditions.

The EDS analysis conducted on the cross-section of the clad layer LC1 produced at the lowest heat input and free cooling, beside the massive eutectic W2C/WC, confirmed presence of M2C type carbides as dendritic precipitations, Figure 15 and Table 5. In turn, the densely packed light grey precipitations between the massive carbides show composition typical for M3C carbides in eutectic, Figure 15, Table 5.

The composition of dark grey matrix indicates presence of Ni/Ni3B eutectic, possibly with some amount of M2C type carbides, Figure 15, Table 5. In general, the dilution was higher from a dozen to several dozen percent in the case of clad layers produced at free cooling conditions, Figure 8. Therefore, the elements from the non-alloy steel substrate, mainly the Fe, passed into the clad. As a result of higher dilution, the composition of the clad is different than the original composition of the applied powder, Table 2.

Although the quantitative results of the measurement of nitrogen content should be considered with a great uncertainty, it is evident that in this case the nitrogen content is significantly lower, on the limit of the measurement error of EDS spectrometer, Table 5. This is due to application of shielding of argon during laser deposition tests, and thus protection against absorption of gases from the ambient. The X-ray diffraction pattern of the experimental powder is presented in Figure 3. It indicates the presence of tungsten carbides δ-WC and β-W2C, as well as the γ-Ni. Additionally, clear picks for Ni3B and complex carbides M7C<sup>3</sup> type were identified, Figure 3.

In turn, the X-ray diffraction patterns for two comparative samples are presented in Figure 16. The comparative samples were prepared by multi-bead cladding of the disc's substrate intended for the tribological tests. The discs were produced by cladding the

substrate with an overlap approx. 20 ÷ 25%, to produce the coating on the entire disc surface. The comparative coatings were produced at the minimum heat input (minimum laser output power) at free cooling and at forced cooling under cryogenic conditions (parameters for the single beads LC1 and HC1, Table 3). Surfaces of the discs were grinded with roughness Ra 0.8 ÷ 1.1 µm. *Materials* **2021**, *14*, x FOR PEER REVIEW 21 of 26

**Figure 16.** XRD patterns of the comparative coatings produced at the same minimum heat input of 60 J/mm (scanning speed 500 mm/min, laser output power 500 W) but different cooling conditions: LC1 at free cooling, and HC1 with forced localised cooling under cryogenic conditions, Table 3. **Figure 16.** XRD patterns of the comparative coatings produced at the same minimum heat input of 60 J/mm (scanning speed 500 mm/min, laser output power 500 W) but different cooling conditions: LC1 at free cooling, and HC1 with forced localised cooling under cryogenic conditions, Table 3.

The initial and general comparison of the two X-ray diffraction patterns show a clear difference in the phase composition. The sample coated at forced cooling under cryogenic conditions (parameters for HC1) shows simply more peaks compared to the sample coated at free cooling (parameters for LC1), Figure 16. However, peaks for the LC1 are more intensive. It is evident that the peaks for γ-Ni are dominant in the case of the coating produced at free cooling (LC1), Figure 16. There are also clear peaks for complex carbides M23C6 and M7C3 type. Moreover, slight peaks for δ-WC and Ni3B were identified. The obtained results are consistent with the previous observations of SEM micrographs and EDS analysis. The phase composition of the clad layer produced under free cooling LC1 is also affected by the higher dilution by the substrate material, mainly iron Fe, as shown the chemical composition analysis. In turn, the X-ray diffraction patterns indicate that the share of eutectic W2C/WC is significantly higher in the coating produced under cryogenic conditions. The dominant γ-Ni phase in the coating produced at free cooling is a result of different thermal conditions of the deposition process, higher impact of the heat, and thus higher degree of carbides decomposition and dissolution, as well as higher dilution by the Fe from the substrate of non-alloy steel. Therefore, the carbides precipitated from the liquid solution are more complex, as shown in Figure 16. The initial and general comparison of the two X-ray diffraction patterns show a clear difference in the phase composition. The sample coated at forced cooling under cryogenic conditions (parameters for HC1) shows simply more peaks compared to the sample coated at free cooling (parameters for LC1), Figure 16. However, peaks for the LC1 are more intensive. It is evident that the peaks for γ-Ni are dominant in the case of the coating produced at free cooling (LC1), Figure 16. There are also clear peaks for complex carbides M23C<sup>6</sup> and M7C<sup>3</sup> type. Moreover, slight peaks for δ-WC and Ni3B were identified. The obtained results are consistent with the previous observations of SEM micrographs and EDS analysis. The phase composition of the clad layer produced under free cooling LC1 is also affected by the higher dilution by the substrate material, mainly iron Fe, as shown the chemical composition analysis. In turn, the X-ray diffraction patterns indicate that the share of eutectic W2C/WC is significantly higher in the coating produced under cryogenic conditions. The dominant γ-Ni phase in the coating produced at free cooling is a result of different thermal conditions of the deposition process, higher impact of the heat, and thus higher degree of carbides decomposition and dissolution, as well as higher dilution by the Fe from the substrate of non-alloy steel. Therefore, the carbides precipitated from the liquid solution are more complex, as shown in Figure 16.

It should be noted that the detection level of the applied XRD method is approx. 3% for the individual phase, therefore, it does not allow for precise identification of the fine precipitations at a low share. However, due to the conjunction with chemical composition analysis in the micro regions and individual constituents, precise detection of the phase composition was provided. It should be noted that the detection level of the applied XRD method is approx. 3% for the individual phase, therefore, it does not allow for precise identification of the fine precipitations at a low share. However, due to the conjunction with chemical composition analysis in the micro regions and individual constituents, precise detection of the phase composition was provided.

#### *3.3. Hardness*

*3.3. Hardness* 

The profiles of hardness were determined on cross-section of the clad layers produced at minimum and maximum heat input at free cooling and under cryogenic conditions. For the Vickers tests of the composite clad layers the applied load was set at 5N. The The profiles of hardness were determined on cross-section of the clad layers produced at minimum and maximum heat input at free cooling and under cryogenic conditions. For the Vickers tests of the composite clad layers the applied load was set at 5N. The

width of the zone with a high hardness. The clad layers produced at free colling show

large load was chosen to avoid large scatter of results (low values for matrix and high for carbides) providing mean hardness value for the subsequent regions. However, the profiles of Vickers HV0.5 hardness presented in Figure 17 showed clear difference for the hardness distribution for individual clad layers. The difference is related mainly with the width of the zone with a high hardness. The clad layers produced at free colling show higher width of the zone with the high hardness ranged approx. from 500HV0.5 to 600 HV0.5. It is related directly with the shape of the individual clads, specifically with the height and penetration depth. The clad layers produced at free cooling conditions have higher penetration depth compared to the clads produced under cryogenic conditions. Therefore, the total depth (width on the graph) of the clad layers produced at free cooling conditions is higher. *Materials* **2021**, *14*, x FOR PEER REVIEW 22 of 26 higher width of the zone with the high hardness ranged approx. from 500HV0.5 to 600 HV0.5. It is related directly with the shape of the individual clads, specifically with the height and penetration depth. The clad layers produced at free cooling conditions have higher penetration depth compared to the clads produced under cryogenic conditions. Therefore, the total depth (width on the graph) of the clad layers produced at free cooling conditions is higher.

**Figure 17**. Hardness distribution on cross-sections of the representative clad layers produced by laser deposition of experimental WC-Ni composite powder (Table 3) at free cooling conditions (LC1, LC4—conventional laser cladding); and at forced localized cooling by nitrogen vapours stream under cryogenic conditions (HC1, HC4—hybrid laser deposition process). **Figure 17.** Hardness distribution on cross-sections of the representative clad layers produced by laser deposition of experimental WC-Ni composite powder (Table 3) at free cooling conditions (LC1, LC4—conventional laser cladding); and at forced localized cooling by nitrogen vapours stream under cryogenic conditions (HC1, HC4—hybrid laser deposition process).

On the other hand, despite the difference in the chemical and phase composition of the clad layers, the maximum values of hardness and the scattering of results are very similar. It should be noted that the presented results of hardness measurement for each point on the graphs are a mean value taken from three different sections for every tested clad layer. The values of standard deviation indicate small dispersion of results. This is due to relatively large size of Vickers's indenter imprint at the load of 5N. The clad layers HC1 produced al lowest heat input and under cryogenic conditions showed the highest value of hardness 520HV0.5 directly under the top surface (face) of the clad. In turn, the highest measured hardness approx. 600HV0.5 was detected for the clad layer produced at maximum het input and cryogenic conditions. For every tested clad layer, starting from a certain depth the hardness decreases gradually till the value of the base metal of nonalloy steel (S235JR) approx. 120 ÷ 15HV0.5. This finding contrasts with the phenomenon described in the previous manuscript on "Laser Deposition of Fe-based metallic powder under cryogenic conditions" [1]. However, different technique of cryogenic cooling was applied in the previous study, consisted of submerging the entire sample in the liquid nitrogen bath. Therefore, the undercooling of the substrate was higher in the previously tested "liquid nitrogen bath" technic. On the other hand, despite the difference in the chemical and phase composition of the clad layers, the maximum values of hardness and the scattering of results are very similar. It should be noted that the presented results of hardness measurement for each point on the graphs are a mean value taken from three different sections for every tested clad layer. The values of standard deviation indicate small dispersion of results. This is due to relatively large size of Vickers's indenter imprint at the load of 5N. The clad layers HC1 produced al lowest heat input and under cryogenic conditions showed the highest value of hardness 520HV0.5 directly under the top surface (face) of the clad. In turn, the highest measured hardness approx. 600HV0.5 was detected for the clad layer produced at maximum het input and cryogenic conditions. For every tested clad layer, starting from a certain depth the hardness decreases gradually till the value of the base metal of non-alloy steel (S235JR) approx. 120 ÷ 15HV0.5. This finding contrasts with the phenomenon described in the previous manuscript on "Laser Deposition of Fe-based metallic powder under cryogenic conditions" [1]. However, different technique of cryogenic cooling was applied in the previous study, consisted of submerging the entire sample in the liquid nitrogen bath. Therefore, the undercooling of the substrate was higher in the previously tested "liquid nitrogen bath" technic.

ditions and different thermal conditions. The values presented at the graphs are mean values taken from at least three individual measurements. The volume loss of the coatings

It must be noted that the task of the study is not a comprehensive analysis of the mechanisms of wear and tribological characteristics. Therefore, just basic results of tribo-

*3.4. Tribological Test* 

#### *3.4. Tribological Test*

It must be noted that the task of the study is not a comprehensive analysis of the mechanisms of wear and tribological characteristics. Therefore, just basic results of tribological ball-on-disc tests were presented for the final comparison of the characteristics and wear resistance of the comparative coatings produced under different technological conditions and different thermal conditions. The values presented at the graphs are mean values taken from at least three individual measurements. The volume loss of the coatings was calculated by determining the wear track profile in four places around the circumference of the wear track. Comparison of the wear tracks for representative samples produced at the lowest heat input but different cooling conditions (coating produced at parameters for HC1 and LC1) is shown in Figure 18. *Materials* **2021**, *14*, x FOR PEER REVIEW 23 of 26 was calculated by determining the wear track profile in four places around the circumference of the wear track. Comparison of the wear tracks for representative samples produced at the lowest heat input but different cooling conditions (coating produced at parameters for HC1 and LC1) is shown in Figure 18.

**Figure 18.** A view of the discs prepared for ball-on-disc wear tests with coating produced by laser deposition; (**a**) coating produced with parameters and technological conditions for the clad HC1 (cryogenic hybrid method), (**b**) coating produced with parameters and technological conditions for the clad LC1 (conventional laser cladding), and after the tribological test; (**c**) sample with coating HC1 type, and (**d**) sample with coating LC1 type. **Figure 18.** A view of the discs prepared for ball-on-disc wear tests with coating produced by laser deposition; (**a**) coating produced with parameters and technological conditions for the clad HC1 (cryogenic hybrid method), (**b**) coating produced with parameters and technological conditions for the clad LC1 (conventional laser cladding), and after the tribological test; (**c**) sample with coating HC1 type, and (**d**) sample with coating LC1 type.

The wear track of the sample HC1 produced under cryogenic conditions and the lowest heat input has the smallest width, as shown in Figure 18c. As expected, the determined volume loss was also the lowest for the coating HC1, Figure 19a. Since the lowest volume loss, the better wear resistance, it means that the composite coating produced at the lowest heat input and cryogenic conditions have the highest wear resistance under dry sliding conditions of the experiment. As can be seen in Figure 19a, even the coating produced at the maximum heat input but at localized forced cooling under cryogenic conditions showed lower volume loos and thus higher wear resistance than the coatings produced at free cooling and the lowest heat input. An explanation for this observed phenomenon can be found in Figure 11a, showing the share of massive carbides in the matrix. The share The wear track of the sample HC1 produced under cryogenic conditions and the lowest heat input has the smallest width, as shown in Figure 18c. As expected, the determined volume loss was also the lowest for the coating HC1, Figure 19a. Since the lowest volume loss, the better wear resistance, it means that the composite coating produced at the lowest heat input and cryogenic conditions have the highest wear resistance under dry sliding conditions of the experiment. As can be seen in Figure 19a, even the coating produced at the maximum heat input but at localized forced cooling under cryogenic conditions showed lower volume loos and thus higher wear resistance than the coatings produced at free cooling and the lowest heat input. An explanation for this observed phenomenon can be found in Figure 11a, showing the share of massive carbides in the matrix. The share

of massive carbides (eutectic W2C/WC) in the matrix of every clad layer produced under cryogenic conditions is higher than for the clad layers produced at free cooling. Even in

2000 W) but with simultaneous localized cooling the share of massive carbides is 22.4%, while in the case of the clad layer produced at the lowest heat input LC1 (500 W) but free cooling the share of carbides is less than 15%. Thus, the results show that under the test conditions, the wear resistance of the composite coatings is proportional to the massive

of massive carbides (eutectic W2C/WC) in the matrix of every clad layer produced under cryogenic conditions is higher than for the clad layers produced at free cooling. Even in the case of the clad layer produced at the maximum heat input HC4 (laser output power 2000 W) but with simultaneous localized cooling the share of massive carbides is 22.4%, while in the case of the clad layer produced at the lowest heat input LC1 (500 W) but free cooling the share of carbides is less than 15%. Thus, the results show that under the test conditions, the wear resistance of the composite coatings is proportional to the massive carbides (eutectic W2C/WC) content in the matrix. On the other hand, the volume loos of the counter face material of steel ball is directly proportional to the volume loos of the discs samples. This phenomenon may be explained by the values of the coefficient of friction determined during individual tests. The values of coefficient of friction presented at the graph in Figure 19b are the mean values, however, a clear relationship between the coefficient of friction and wear intensity can be observed. In general, the higher coefficient of friction the higher wear intensity, both the sample and the steel ball, Figure 19. The obtained results indicate that the coefficient of friction also depends on the phase composition of the coatings, mainly on the share of massive carbides in the matrix, as well as composition of the matrix. *Materials* **2021**, *14*, x FOR PEER REVIEW 24 of 26 carbides (eutectic W2C/WC) content in the matrix. On the other hand, the volume loos of the counter face material of steel ball is directly proportional to the volume loos of the discs samples. This phenomenon may be explained by the values of the coefficient of friction determined during individual tests. The values of coefficient of friction presented at the graph in Figure 19b are the mean values, however, a clear relationship between the coefficient of friction and wear intensity can be observed. In general, the higher coefficient of friction the higher wear intensity, both the sample and the steel ball, Figure 19. The obtained results indicate that the coefficient of friction also depends on the phase composition of the coatings, mainly on the share of massive carbides in the matrix, as well as composition of the matrix.

**Figure 19.** Volume loss of the counter body steel ball and the disc specimens of commercially pure titanium, titanium alloy and the nitrided surface layers and after sliding for a distance of 188.4 m at normal load of 30 N (**a**), and comparison of the coefficient of friction of the disc specimens sliding with a steel ball during ball-on-disc wear tests (**b**). **Figure 19.** Volume loss of the counter body steel ball and the disc specimens of commercially pure titanium, titanium alloy and the nitrided surface layers and after sliding for a distance of 188.4 m at normal load of 30 N (**a**), and comparison of the coefficient of friction of the disc specimens sliding with a steel ball during ball-on-disc wear tests (**b**).

#### **4. Conclusions 4. Conclusions**

The novel technique of laser deposition of composite WC-Ni powder with forced localized cooling of the deposit by nitrogen vapours stream under cryogenic conditions was successfully demonstrated and its potential for shaping the geometry and the microstruc-The novel technique of laser deposition of composite WC-Ni powder with forced localized cooling of the deposit by nitrogen vapours stream under cryogenic conditions was successfully demonstrated and its potential for shaping the geometry and the mi-

ture was presented. The results of the study confirmed the advantageous effect of the forced cooling in terms of chemical and phase composition, microstructure, and wear re-

reduces the tendency to decomposition and to dissolve carbides in the melt pool during laser deposition of composite WC/Ni powder. The share of carbides on cross section of the clad layer produced at minimum heat input (laser output power 500 W and scanning speed 500 mm/min) at conventional conditions of free cooling was less than 15%, while

crostructure was presented. The results of the study confirmed the advantageous effect of the forced cooling in terms of chemical and phase composition, microstructure, and wear resistance under dry sliding conditions. The use of the forced cooling technique significantly reduces the tendency to decomposition and to dissolve carbides in the melt pool during laser deposition of composite WC/Ni powder. The share of carbides on cross section of the clad layer produced at minimum heat input (laser output power 500 W and scanning speed 500 mm/min) at conventional conditions of free cooling was less than 15%, while thanks to the application of forced localized cooling of the deposit the share of carbides was maintained over 51%. Moreover, the degree of dilution was significantly reduced thanks to the forced localized cooling. The chemical composition of clad layer produced at minimum heat input and with forced cooling under cryogenic conditions was very similar to that of experimental composite WC-Ni powder, except for the increased nitrogen content in the dendritic matrix and interdendritic regions. The increased content of nitrogen detected in the matrix of the clad layer is resulted by presence of high amount of gaseous nitrogen in deposition region if the nitrogen vapours stream is applied. Therefore, during the laser deposition of composite powder (laser cladding) the alloying process takes place additionally. The material of the clad layer (the deposit) is simultaneously enriched by nitrogen from the gaseous atmosphere, so the demonstrated process of laser deposition of composite powder with simultaneous localized cooling by the nitrogen vapor stream can be considered as hybrid process, combining the conventional laser cladding (LC) and also laser gas nitriding (LGN).

**Author Contributions:** Conceptualization, methodology, investigation, formal analysis, validation, supervision, writing—review and editing, A.L.; project administration, funding acquisition, D.S. All ´ authors have read and agreed to the published version of the manuscript.

**Funding:** The research was supported by the National Centre for Research and Development, Poland, under the grant number POIR.01.01.01-00-0278/15-007-02, financed by EU funds.

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** The data presented in this study are available on request from the corresponding author. The data are not publicly available due to know how protection.

**Acknowledgments:** The authors thank for technical support to all those who assisted in the preparation of samples and the conducting of measurements, especially to Marcin Zuk, Adrian Kukofka, ˙ and students who assisted in the laboratory tests.

**Conflicts of Interest:** Authors declare no conflict of interest.

#### **References**

