Abstract

In recent years, cold spray additive manufacturing (CSAM) has become an attractive technology for surface modification and protection. However, due to the intrinsic porous nature of CSAM coatings, they suffer from rapid material degradation due to premature brittle fracturing induced by tribological interactions. In this work, laser shock peening (LSP) was utilized as a post-processing technology to mitigate the surface porosity and augment the surface characteristics of CSAM 316L stainless steel (SS). Due to the synergistic influence of severe plastic deformation and rapid surface heating, the surface porosities were effectively healed, thus reducing the surface roughness. Combined with the surface-strengthening effects of LSP, the frictional resistance and transfer layer formation on the CSAM LSP surfaces were reduced. The underlying mechanisms for these findings were discussed by correlating the atomic, microstructural, and physical features of the LSP surfaces. Based on these findings, it can be suggested that LSP is indeed a useful technique to control the surface characteristics of CSAM 316L SS coatings.

1 Introduction

For many industrialized components, the presence of mechanical abrasion is quite common as they tend to be exposed to a wide array of working environments [1]. This is especially true for standard engineering components such as gears, bearings, pistons, injectors, and metal forming/cutting tools where their surfaces are continually in contact with other surrounding components. The plastic deformation over time will lead to eventual material degradation thus shortening the in-service lifetime of these components [2]. For example, in applications such as nuclear power generation, minimizing friction and wear can reduce the energy consumption of various processes which can improve their efficiencies and overall reliability [3]. Extensive research efforts have been paid to achieve this goal. First, the material of use can be improved either by (1) the replacement of more durable material, (2) employing some type of composite during the fabrication process, or simply by (3) changing the fabrication process altogether. However, these methods suffer from high costs and low feasibility for large-scale applications [4]. In place of this, surface processing technologies have attracted much attention as an alternative method of friction and wear control [511].

Among a variety of surface processing technologies, cold spray additive manufacturing (CSAM) is one surface enhancement technique that has seized the attention of many due to its unique solid-state process [12,13]. As a relatively novel technique, CSAM is essentially a low-temperature process that depends on the kinetic energies of impacted particles for material deposition [1416]. This is quite advantageous in the aspect that the original phase of the particle feedstock is preserved in contrast to traditional thermal-based techniques such as high-velocity oxy-fuel (HVOF) spray and plasma spraying [17]. In its most basic sense, CSAM consists of rapidly launching micrometer-sized particles (typically within the range of 300–1200 m/s [18]) to a working substrate through a de Laval nozzle at supersonic velocities [19]. For these speeds to be achieved, high-pressure gas is applied, typically in the form of helium, nitrogen, or air, which is typically categorized as high-pressure cold spraying (HPCS) [20]. As the particles are rapidly impacted, mechanical bonding is induced by the combination of shear stresses from both the substrate and the coating [21]. Compared to other thermal-spray technologies, there is no particle melting, which will preserve undesirable defects such as grain diffusion, phase transformations, and tensile residual stresses [19].

Given the simplicity and ease of application for CSAM coatings, many industrial sectors such as thermal power plants, nuclear power plants, automotive, aerospace, and marine have greatly benefitted from the prolonged lifespans of their respective components [22,23]. Despite its wide applications, there are some intrinsic disadvantages associated with the CSAM process which has limited its full potential in improving its wear resistance. The surface defects formed during the CSAM process, such as porosity, are quite detrimental to the tribological properties of the coatings [24]. This defect is mainly due to the lack of heterogeneous deformation of the individual particles with respect to the impacted region. Whether it is from the irregularity of the particle morphology/size or from the lack of sufficient plastic deformation, microcracks, and micropores are typically observed within the inter-particle boundaries which can limit the structural integrity of the coating [25]. This in combination with the already brittle nature of these coatings creates early crack propagation which can increase the wear-rate and tendency of third-body wear in the affected region [26]. As consequence, the adherence of cold spray (CS) coatings is negatively impacted by the unmatched tensile and residual stresses at both the top and bottom of the coating which then will further promote early crack initiation [27].

Post-surface processing techniques have been explored as a practice routine to tailor and optimize the tribological properties of CSAM coatings. Although many have studied the effects of post-processing techniques such as heat treatment [28], friction stir processing [29], and shot peening [30] on CSAM coatings, there are few publications that study the effects of laser shock peening (LSP) on CSAM coatings [31,32], especially in the sense of pore reduction. LSP is a laser-based surface processing technique that has been widely utilized to enhance the tribological properties of metallic materials [3337]. During this process, an intense laser pulse with a duration time of 1–30 ns is typically utilized to generate high shockwave pressure (in an order of GPa) through laser-matter interactions [38]. Throughout this process, a transparent confinement and an ablative coating are typically utilized to absorb the laser energy as well as create a constrained plasma to the surface. However, it should be mentioned that in some cases the transparent confinement and/or ablative coating can be removed during peening [39]. As the laser-induced shockwave propagates through the specimen, the processed material undergoes ultrahigh strain rates (106–107/s) plastic deformation, leading to localized work hardening and grain refinement effect. Additionally, deep compressive stresses are introduced to the material which can enhance the fatigue performance of the sample processed due to its capability to balance external cyclic stresses [40]. From an application standpoint, this will limit the likelihood of early crack propagation which can result in extended operational lifespans of various components [4043].

Due to its advantages, LSP has been applied as a post-processing technique for a wide variety of materials and applications. Compared to technologies such as shot peening or thermal post-processing techniques (such as laser cladding), LSP can maintain a superior surface quality while preventing the formation of tensile stresses along the surface. This has especially been noted for laser-based additively manufactured (AM) materials that tend to suffer from unwanted pores and voids [44,45]. With many thoroughly studying the effects of LSP on laser-based AM materials, rare research efforts have been paid to investigate the effects of LSP on CSAM coatings. Given its capability to hinder crack propagation and tailor surface roughness, LSP may serve as an efficient approach for CS post-processing. To fill this gap, the feasibility of LSP on CSAM 316L stainless steel (SS) is investigated. Particularly, the surface characteristics and tribological performance will be evaluated. To characterize these changes, a series of experiments were performed on each specimen. These tests include single scratch ball-on-flat tests, microhardness tests, and X-ray diffraction (XRD). The surface morphology of the tested surfaces was also characterized through a scanning electron microscope (SEM) and optical profilometry in order to obtain a better understanding of the relationship between surface morphology and tribological performance.

2 Materials and Methods

2.1 Materials.

In this work, the CSAM deposition process was done using an Impact Spray 5/11 CS system (Impact Innovation GmbH, Rattenkirchen, Germany). For the working material, commercially pure (CP) 316L SS powder was employed as a freestanding deposit. From this point onward, the CSAM substrate will be referred to as the CS substrate. A schematic illustration of the operating system is shown in Fig. 1 [15].

Fig. 1
A schematic illustration of the CSAM process
Fig. 1
A schematic illustration of the CSAM process
Close modal

The particle size prior to deposition was measured to be 20.17 μm with a standard deviation of 6.64 μm. Micrographs used for the particle measurements as well as a bin distribution of the particle size are shown in Fig. 2 It can be seen that the particles are mainly spherical with the exception of some outlying particles. In the case of non-spherical particles, they may lead to uneven stress concentrations upon impact thus leading to porosity [46]. All other processing parameters for the CSAM deposition process are depicted in Table 1.

Fig. 2
Micrograph of the 316L powder feedstock used in this work, scanned from a SEM at (a) 1000× magnification as well as the (b) powder feedstock size distribution using a bin analysis
Fig. 2
Micrograph of the 316L powder feedstock used in this work, scanned from a SEM at (a) 1000× magnification as well as the (b) powder feedstock size distribution using a bin analysis
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Table 1

CSAM Process parameters

Propellant gasSprayed materialGas temperature (C°)Gas pressure (MPa)Spray angle (deg)Stand-off distance (mm)Step size (mm)Powder feed rate (RPM)Powder carrier gas flowrate (m3/h)Nozzle type
N2316L850–11004.0–5.090 deg25.40.7622.0–3.03.0–4.0SiC water-cooled nozzle
Propellant gasSprayed materialGas temperature (C°)Gas pressure (MPa)Spray angle (deg)Stand-off distance (mm)Step size (mm)Powder feed rate (RPM)Powder carrier gas flowrate (m3/h)Nozzle type
N2316L850–11004.0–5.090 deg25.40.7622.0–3.03.0–4.0SiC water-cooled nozzle

2.2 Laser Shock Peening.

Post-deposition, the CS coating was sectioned into four pieces for LSP and polished using traditional sandpaper polishing to a grit size of 1200. Final polishing was performed using a 0.05 μm low-viscosity alumina-slurry paste to achieve a uniform mirror surface finish with a surface roughness (Sa) of around 0.05 ± 0.03 μm. Afterward, LSP was performed using a Q-switched Nd-YAG laser (Surelite III, Continuum, Inc., Miramar Beach, FL, USA) at an operating wavelength of 1064 nm using a 50% overlap ratio. In order to achieve this overlap rate, a stepper motor linear translation stage (NRT100, ThorLabs, Newton, NJ, USA) was utilized to control the movement of the sample during LSP processing. For better visualization, the LSP setup and peening path are shown in Fig. 3. During peening, a pulse width of 5 ns at the full width at half maxima was employed. The laser beam diameter was measured to be 2.0 mm. Black vinyl tape with a thickness of 130 μm was utilized as the ablative coating whereas no confinement media was used in this work. The usage of no confinement media has been reported to be an acceptable technique that can achieve similar surface-strengthening effects as LSP with a confinement media [4749]. Additionally, it can be speculated that the thermal effects of LSP without confinement media can potentially seal the porosity of porous surfaces. By varying the Q-switch delay time, the laser intensities can be adjusted and their values can be calculated using the following formula [37]:
I=EPW×A
(1)
where I is the laser density (GW cm−2), E is the pulse energy (J), PW is the pulse width (ns), and A is the laser spot area (cm2). The calculated laser intensities in the present study are shown in Table 2. From this point onward, the peened CS substrates will be referred to as CS + I1, CS + I2, CS + I3, and CS + I4. The increase in sample numbering will represent the laser intensity applied, consisting of 1.91 GW cm−2 (CS + I1), 2.54 GW cm−2 (CS + I2), 3.18 GW cm−2 (CS + I3), and 3.81 GW cm2 (CS + I4), as shown in Table 2.
Fig. 3
(a) A visual view of the LSP equipment used in this work as well as the (b) corresponding path used for peening
Fig. 3
(a) A visual view of the LSP equipment used in this work as well as the (b) corresponding path used for peening
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Table 2

LSP Parameters used for the CS specimens

SpecimenLaser intensity, I (GW cm−2)
Sample #1 (CS)0
Sample #2 (CS + I1)1.91
Sample #3 (CS + I2)2.54
Sample #4 (CS + I3)3.18
Sample #5 (CS + I4)3.81
SpecimenLaser intensity, I (GW cm−2)
Sample #1 (CS)0
Sample #2 (CS + I1)1.91
Sample #3 (CS + I2)2.54
Sample #4 (CS + I3)3.18
Sample #5 (CS + I4)3.81
Post-processing, the volumetric strain induced by the shockwave peak pressure was additionally calculated using the well-reputed shockwave pressure model created by Fabbro [50] which is expressed by
P=0.01(α2α+3)×Z×I
(2)
where I is the calculated laser intensity (GW cm−2), α is the interaction efficiency (which is typically calculated at 0.15 [51]), Z is the shockwave impedance of the material, and P is the laser shock wave pressure (GPa). To solve for Z, the following equation was employed:
Z=2(1Z1)+(1Z2)
(3)
where the acoustic impedance of the material and confinement media is Z1 and Z2. According to the above formula, the peak shockwave pressure for CS + I1, CS + I2, CS + I3, and CS + I4 was found to be 4.41 GPa, 7.21 GPa, 9.19 GPa, and 10.79 GPa, as shown in Fig. 4.
Fig. 4
The calculated shockwave pressure of the laser intensities used in this work
Fig. 4
The calculated shockwave pressure of the laser intensities used in this work
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2.3 Characterization Techniques.

The surface morphologies of the CS coatings processed before and after LSP were characterized using a Rtec 3D optical profilometer (Rtec-Instruments, San Jose, CA, USA) with a resolution of 50 nm. A surface area of 1.9 × 1.2 mm2 was used to determine the change of surface roughness (SR) parameters induced by LSP. These surface parameters were calculated using the Gwyddion software [52]. To further characterize the change in surface morphology, micrographs of the surface porosity and microstructure were obtained from a JSM-6010LA InTouchScope SEM (JEOL, Tokyo, Japan) for the surfaces of the peened and non-peened specimens. Before microstructural evaluation, all samples were submerged in Carpenter’s Etchant (8.5 g FeCl3 + 2.4 g CuCl2 + 122 ml C2H5OH + 122 ml HCl + 6 ml HNO3) [53]. To investigate the possible phase transformation during the LSP process, a Bruker-D2 Phaser (Bruker, Madison, WI, USA) XRD apparatus with Cu-Kα radiation was employed. The excitation voltage and current were set to 40 kV and 25 mA. The 2θ angles spanned from 20 deg to 100 deg with a scan speed of 1 deg/min. The step width was set to 0.02 deg throughout each scan. Surface microhardness (HV) values also were measured post-processing in order to further obtain a better understanding of the surface-strengthening effect of LSP. These readings were measured under the loading conditions of 100-g with a 10-s dwell time (Beuhler-Wilson, Tukon 1202). Each test was replicated ten times to ensure repeatability.

2.4 Tribological Testing.

For tribological evaluation, each sample was subjected to a dry single scratch test as per ASTM G133 [54]. These tests were performed using a Rtec multi-function Tribometer (Rtec-Instruments, San Jose, CA, USA) at an ambient room temperature of ∼25 °C. The counter material used was an Al 6061-T6 pin. For each test, a normal load of 10 N was applied for a sliding distance of 10 mm. The moving velocity was set to 0.5 mm s−1. Afterward, the transfer layer formation of the Al pin on the CS flat was studied using an SEM and energy dispersive spectroscopy (EDS) from a JSM-7100FT SEM (JEOL, Tokyo Japan). The transfer layer was then mapped and analyzed. To ensure repeatability, each test was replicated three times.

3 Results

3.1 Surface Roughness and Morphology.

To understand the influence of LSP on the SR of the CS samples, the change in average surface roughness (Sa) was investigated. The variations of this parameter with respect to laser intensity are shown in Fig. 5. As expected, the influence of LSP at 1.91 GW cm−2 (i.e., CS + I1) roughened the CS surface by 17.9%. This is to be expected due to the peak pressure shockwave inducing plastic deformation to the surface [55]. However, interestingly, Sa begins to decrease after being subjected to a laser intensity of 2.54 GW cm−2 (i.e., CS + I2). In fact, relative to the unpeened specimen, the Sa value for the CS + I2 specimen decreased by 4.7%. Upon increasing laser intensity, the CS + I3 and CS + I4 specimens also showed a reduction in Sa by 30.9% and 39.5%.

Fig. 5
The change in Sa of the CS, CS + I1, CS + I2, CS + I3, and CS + I4 specimens
Fig. 5
The change in Sa of the CS, CS + I1, CS + I2, CS + I3, and CS + I4 specimens
Close modal

To obtain additional information regarding the change in surface characteristics with respect to laser intensity, 3D surface profiles alongside their corresponding 2D line profile characterization were effectively analyzed (Fig. 6). It can be seen that across all specimens, there was some presence of surface porosity, as indicated by the listed arrows. This is to be expected due to the intrinsic defects of the CS process, as aforementioned. However, as the laser intensity increased, the surface asperities were effectively deformed and reconstructed, thus sealing most of the surface porosity as a function of laser intensity. This finding is not uncommon, as LSP has been reported to seal surface defects due to the compressive stresses induced by peening [5660].

Fig. 6
The change in surface morphology for samples (a–b) CS, (c–d) CS + I1, (e–f) CS + I2, (g–h) CS + I3, and (i–j) CS + I4. The arrows indicate the regions of porosity.
Fig. 6
The change in surface morphology for samples (a–b) CS, (c–d) CS + I1, (e–f) CS + I2, (g–h) CS + I3, and (i–j) CS + I4. The arrows indicate the regions of porosity.
Close modal

3.1.1 Microstructure.

In addition to the change in surface morphology, LSP is widely known for its capability of surface grain refinement due to the ultrahigh strain rate-induced dislocation dynamics [61]. Figure 7 depicts the change in microstructure across all specimens. Observing the CS substrate (Fig. 7(a)), the various impacted particles can be clearly differentiated from one another by the black lines. The lengthened morphology of the particles clearly indicates that mechanical and metallurgical bonding (due to the translation of kinetic energy of the rapidly accelerated particles) did occur due to the severe plastic deformation from the deposition process [20]. Post-peening, there is a remarkable distinction between the CS substrate’s structural morphology compared to its original pre-peened state. Across all peened samples (Figs. 7(b)7(e)), non-uniform deformation bands are densely packed within the grain boundaries with minimal spacing. Upon closer inspection, the microstructure of the CS specimens at lower intensities appears to be largely cellular with homogeneous nanoscale grains. However, at higher laser intensities, the cellular structures appear to be increasingly elongated likely due to the rapid melting from the laser peening. Nonetheless, the microstructural features of all peened specimens are visibly refined from their original state due to the severe plastic deformation from the LSP process.

Fig. 7
The etched surface microstructure of the (a) CS, (b) CS + I1, (c) CS + I2, (d) CS + I3, and (e) CS + I4 specimens
Fig. 7
The etched surface microstructure of the (a) CS, (b) CS + I1, (c) CS + I2, (d) CS + I3, and (e) CS + I4 specimens
Close modal

It can also be noted that the degree of surface porosity also decreased as laser treatment was applied to the specimens. This can be visually seen from the surface micrographs that show the surfaces of the substrates alongside their pore area calculations from the ImageJ software in Fig. 8. For the CS substrate (Figs. 8(a) and 8(b)), various pores can be seen along the surface. However, considering that no confinement media was used, the plasma formation from the laser peen was not confined, which likely resulted in the surface experiencing a form of atomistic diffusion [6264]. As such, the localized particles are reorganized into a more robust and dense structure as a function of laser intensity, as shown in Figs. 8(c)8(j) [65,66]. A quantitative analysis of the pore count can be shown in Fig. 9, which confirms that the percentage of porosity decreases due to LSP treatment. For the CS substrate, a density of 99.15% was recorded, whereas the CS + I1, CS + I2, CS + I3, and CS + I4 substrates had densities of 99.35%, 99.54%, 99.73%, and 99.92%.

Fig. 8
The porosity morphology of the (a) CS, (b) CS + I1, (c) CS + I2, (d) CS + I3, and (e) CS + I4 specimens
Fig. 8
The porosity morphology of the (a) CS, (b) CS + I1, (c) CS + I2, (d) CS + I3, and (e) CS + I4 specimens
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Fig. 9
The change in porosity for the CS and CS + LSP specimens
Fig. 9
The change in porosity for the CS and CS + LSP specimens
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From a cross-sectional perspective, micrographs of the CS specimen (without peening), CS + I2, and CS + I4 specimens are shown in Fig. 10. It can be seen that for the CS specimen (Figs. 10(a)10(d), there appears to be a large number of pores dispersed along the cross section. However, when subjected to peening, the porosity along the surface (Figs. 10(b), 10(c), 10(e) and 10(f)) greatly decreases. This is especially evident for the CS + I4 specimen (Figs. 10(c)10(f)), of which the depth of the pore closure appears to be the greatest when compared to the CS + I2 specimen (Figs. 10(b)10(e)). To explicate, the surface recoil pressure and vaporization of the plasma were effectively confined in the ablative coating, which likely assisted with surface densification. As the depth increases, the extreme plastic deformation induced by the peak shockwave pressure allowed for compressive stresses to seal the pores, thus allowing for the affected region to be effectively densified. This finding is supported by Glaser et al. [67] as well as other peening-related literature [68].

Fig. 10
The change in the whole cross-sectional porosity as well as magnified upper cross-sectional porosity for the (a,d) CS, (b,e) CS + I2, and (c,f) CS + I4 specimens
Fig. 10
The change in the whole cross-sectional porosity as well as magnified upper cross-sectional porosity for the (a,d) CS, (b,e) CS + I2, and (c,f) CS + I4 specimens
Close modal

3.1.2 Phase Analysis.

To further quantify the degree of plastic deformation from LSP, XRD patterns of the pre-and-post-peened specimens were investigated, as shown in Fig. 11. Based on the XRD spectra, all samples exhibit a gamma (γ) structure with no phase changes induced by LSP. Typically in LSP literature, a martensite transformation occurs due to the high degree of plastic strain from peening. According to Chen et al. [69], strain rates as low as 101 and 10−3 s−1 induced martensite (ε) transformations in stainless steels due to the high degree of dislocation motions. With LSP-inducing strain rates as high as 106 s−1 [70], it would be expected that some degree of a phase transformation would occur. However, the localized surface heating from the laser irradiation process could potentially suppress the γ → ε transformation, as the degree of plastic strain is suppressed [71]. Additionally, due to the strain rates of LSP being much higher than the necessary strain rates for a martensite transformation, the excessive amount of dislocation twins can impede dislocation activation thus impeding the formation of martensite [72].

Fig. 11
XRD Spectra of the CS and CS + LSP specimens
Fig. 11
XRD Spectra of the CS and CS + LSP specimens
Close modal

Despite no phase change, the full width at half maxima (FWHM) of each peak was observed to increase, as shown in Fig. 12(a). It is well known that when a peak is broadened, it indicates that the microstructure of the material is both refined and re-oriented due to the presence of external stresses, which follows suit with the influence of surface peening [73]. Understanding this, the widely utilized Williamson–Hall method [74,75] can be applied to determine the crystallite size, micro-strain, and dislocation density of the specimens. These findings are shown in Figs. 12(b)12(d) and Table 3. According to these findings, the crystallite size for CS + LSP specimens (Fig. 12(b)) showed a net decrease compared to the unpeened CS specimen. Interestingly enough, the CS + I1 specimen showed the greatest reduction in crystallite size (with a 52.9% reduction), with each following laser intensity resulting in a gradual increase. This likely is a result of the microstructural re-orientation observed in Fig. 7, in which the microstructures at higher intensities appear to be less uniform/refined and more elongated. Also, the intensity of the laser at higher laser intensities likely resulted in greater diffusion, which inhibited the formation of finer crystallite sizes. A similar trend can also be seen with dislocation density (Fig. 12(d)), as crystallite size tends to share an inverse relationship to dislocation accumulation. Nonetheless, across all specimens, the crystallite size and dislocation density yielded improvements over the base CS substrate. For the strain (Fig. 12(c)), there was an initial decrease with a gradual increase as a function of laser intensity. Although it might be counterintuitive, the initial decrease in strain can be attributed to the initial microstructural reorientation (due to diffusion) observed in Fig. 7 [76], as CS substrates in their base form tend to already be extremely work-hardened. However, as the laser intensity increases, the degree of plastic deformation due to the higher strain rates (in addition to the rapid heating) occurs resulting in greater flow stresses. With the following conditions, the strain is increased despite the elongated microstructural evolution [77].

Fig. 12
(a) FWHM and corresponding (b) crystallite size, (c) strain, and (d) dislocation density of the CS and CS + LSP specimens
Fig. 12
(a) FWHM and corresponding (b) crystallite size, (c) strain, and (d) dislocation density of the CS and CS + LSP specimens
Close modal
Table 3

The crystallite size, dislocation density, and strain of the CS and CS + LSP specimens calculated from the W–H analysis

Laser intensity (GW cm−2)Crystallite size (nm)Micro-strain (ɛ)Dislocation density (nm−2)
080.43 ± 3.462.60 × 10−3 ± 6.36 × 10−57.42 × 10−4 ± 1.38 × 10−5
1.9137.90 ± 4.032.15 × 10−3 ± 2.01 × 10−41.30 × 10−3 ± 1.70 × 10−5
2.5448.84 ± 1.662.53 × 10−3 ± 8.06 × 10−51.19 × 10−3 ± 7.52 × 10−5
3.1859.38 ± 5.552.80 × 10−3 ± 9.90 × 10−51.09 × 10−3 ± 6.31 × 10−5
3.8161.58 ± 5.392.90 × 10−3 ± 9.19 × 10−51.08 × 10−3 ± 6.05 × 10−5
Laser intensity (GW cm−2)Crystallite size (nm)Micro-strain (ɛ)Dislocation density (nm−2)
080.43 ± 3.462.60 × 10−3 ± 6.36 × 10−57.42 × 10−4 ± 1.38 × 10−5
1.9137.90 ± 4.032.15 × 10−3 ± 2.01 × 10−41.30 × 10−3 ± 1.70 × 10−5
2.5448.84 ± 1.662.53 × 10−3 ± 8.06 × 10−51.19 × 10−3 ± 7.52 × 10−5
3.1859.38 ± 5.552.80 × 10−3 ± 9.90 × 10−51.09 × 10−3 ± 6.31 × 10−5
3.8161.58 ± 5.392.90 × 10−3 ± 9.19 × 10−51.08 × 10−3 ± 6.05 × 10−5

3.1.3 Microhardness.

Figure 13 depicts the change in microhardness of the pre- and-post-peened specimens from both a surface (Fig. 13(a)) and cross-sectional (Fig. 13(b)) point of view. It can be seen that in both cases, LSP does have a positive influence on surface hardness. As shown in Fig. 13(a), prior to LSP, the CS substrate had a surface microhardness value of 345.8 ± 7.1 HV0.1. Post-LSP, the microhardness of the CS substrate gradually increased, with the greatest value being at 427.06 ± 7.99 HV0.1 (with a 23.5% net increase) for the CS + I4 specimen. The cross-sectional hardness readings (Fig. 12(b)) near the peened surface are also in agreement with the surface hardness measurements. At around 1.2 mm, the cross-sectional hardness for the CS + LSP substrates reached equilibrium. Others have also seen similar hardening depths [78]. These results are to be expected as per the Hall–Petch relationship, which can be described as [79,80]:
H=Ho+kHd12
(4)
where d is the grain size, and HokH are relating constants to the hardness of the substrate. Due to the dynamic recrystallization of the peening process, the refinement in crystallinity can be one factor that increased the hardness. However, other factors such as the change in surface morphology (i.e., porosity) also likely influenced the hardness readings. According to Cherry et al. [81], the reduction of surface porosity from laser processing can increase the hardness. It was explained that with porous surfaces, the presence of external loading (from the hardness indenter) along the pore boundaries can result in the local region collapsing thus creating a greater degree of plastic deformation. However, with denser surfaces (such as in this work), there is greater resistance to external loadings.
Fig. 13
(a) Surface and (b) cross-sectional microhardness of the CS and CS + LSP specimens
Fig. 13
(a) Surface and (b) cross-sectional microhardness of the CS and CS + LSP specimens
Close modal

3.2 Tribological Performance

3.2.1 Friction.

When subjected to tribological loading, the effect of LSP drastically reduced the coefficient of friction (COF) as demonstrated in single scratch tests under dry conditions, as shown in Fig. 14(a). For all specimens, the COF was relatively stable throughout sliding. When subjected to LSP, there were some slight oscillations with the COF during sliding. This was especially evident for specimen CS + I1. However, upon increasing the laser intensity, a more-or-less monotonic behavior was observed after the rapid run-in period, which was largely pronounced within the first few seconds of sliding. This stabilization can also indicate that there is no significant plastic deformation occurring along the surface, thus suggesting less tangential stresses along the tribo-loaded interfaces [82,83]. When considering their averaged values (Fig. 14(b)), the COF does indeed follow a direct decrease as a function of laser intensity. The reduction was quantified by 53.6% for the CS + I1 specimen and 60.6%, 65.5%, and 72.6% for the CS + I2, CS + I3, and CS + I4 specimens.

Fig. 14
Change of COF for the CS and CS + LSP specimens
Fig. 14
Change of COF for the CS and CS + LSP specimens
Close modal

3.2.2 Transfer Layer.

Observing the wear track of the specimens (Figs. 15(a), 15(c), 15(e), 15(g), and 15(i)), EDS observations (Figs. 15(b), 15(d), 15(f), 15(h), and 15(j)) were made in order to quantify the degree of Al-6061 transfer material that occurred during tribo-loading. Following a similar trend as the aforementioned COF, there was a gradual decrease of transferred Al to the CS surface as the laser intensity increased. However, this was visually not the case for sample CS + I1 (Fig. 15(d)), as there appears to be slightly more Al accumulation than in sample CS + I0 (Fig. 15(b)). Specifically, the CS + I1 specimen had a 5.33% transfer layer, whereas the CS substrate had a 4.61% transfer layer. In the case of the CS + I2 (Fig. 15(f)), CS + I3 (Fig. 15(h)), and CS + I4 (Fig. 15(j)) specimens, the transfer layer concentration, and disbursement were largely decreased by 2.78%, 1.38%, and 0.33%. SEM observations (Figs. 15(a), 15(c), 15(e), 15(g), and 15(i)) also indicate that the presence of the wear track becomes increasingly faint as the laser intensity increases. This finding insinuates that a lesser degree of plastic deformation occurred during tribo-loading, which further supports the aforementioned friction findings. Likely, the smoothening of the surface asperities allows for greater resistance to shearing from the counter interface, which resulted in a wear track largely free from visual deformation. To quantify this change, the width of the wear tracks was measured in order to further understand the change in tribological response across the specimens. The wear track width for the CS, CS + I1, CS + I2, CS + I3, and CS + I4 specimens was recorded at 149.84 ± 12.20 μm, 120.79 ± 10.41 μm, 102.63 ± 4.36 μm, 101.67 ± 4.62 μm, and 95.24 ± 2.81 μm. This change was measured against the change in aluminum transfer layer content, as shown in Fig. 16. It can be seen that the influence of LSP allows for the wear track width and transfer layer content to share the same decreasing trend, which is again attributed to the smoothening and strengthening of the CS surface post-LSP.

Fig. 15
SEM scans (left-hand side) and EDS maps (right-hand side) of the wear tracks for the (a–b) CS, (c–d) CS + I1, (e–f) CS + I2, (g–h) CS + I3, and (i–j) CS + I4 specimens
Fig. 15
SEM scans (left-hand side) and EDS maps (right-hand side) of the wear tracks for the (a–b) CS, (c–d) CS + I1, (e–f) CS + I2, (g–h) CS + I3, and (i–j) CS + I4 specimens
Close modal
Fig. 16
The aluminum transfer layer % and wear track width of the CS and CS + LSP specimens
Fig. 16
The aluminum transfer layer % and wear track width of the CS and CS + LSP specimens
Close modal

4 Discussion on Tribological Mechanisms

To understand the tribological mechanisms, three synergistic factors should be taken into consideration. The first factor lies in the response of the surface morphology to tribo-loaded deformation, whereas the second and third factors lie in the material-related properties and surface texture of the tribo-loaded regions. To elucidate the first factor, it is important to keep in mind the reduction in surface porosity due to the laser treatment (Fig. 7). Under tribo-loading, it is well known that the influence of pores can increase the contact area of the counter material (i.e., Al6061 pin) [53,84,85]. This creates an implication that the degree of deformation along a porous edge would increase, thus creating greater frictional forces compared to a non-porous surface [86]. Since the CS steel substrate also has a higher hardness than the Al6061 pin, this idea is exacerbated, as surface defects (in the form of porosity) can result in a great degree of plowing on the counter surface, thus increasing the localized frictional force [87]. From a material transfer point of view, the surface pores can also act as active sites that would actively accumulate the counter Al material surface. The accumulation of material in the pore sites was quite evident in the EDS maps (Fig. 15), where the distribution of Al along the wear track generally becomes less scattered as the laser intensity increases. This would suggest that the surface pores do indeed work to accumulate the counter Al material to the surface. It should be mentioned that in the case of sample CS + I1, there is a slightly greater degree of Al accumulation than in the CS specimen. To explicate, the surface features should be considered. In the case of sample CS + I1, the size of the pores is somewhat similar to the CS specimen (Fig. 8). With sample CS + I1 also having a greater surface roughness, the degree of surface shearing (and frictional force) to overcome the roughened surface is increased. As such, the plowing forces on the surface-pin interface increase which would result in a greater transfer layer [88]. From another perspective, the porous regions can also be sensitive to elevated stress concentrations, which could result in brittle fracturing thus extending the exposed porous region and resulting in greater Al accumulation [8991].

The change in material properties also has a large influence on the tribological response of the LSP specimens. This was especially evident with the microstructural features of all peened specimens, as grain refinement due to the ultra-high strain rate of plastic deformation from the LSP allowed for the formation of a denser microstructure. The crystallite size and dislocation density measurements also confirm this finding. According to the findings, the role of crystallite size in the tribological characteristics of the peened specimens was more dominant with lower intensities, whereas the dominant factor at higher intensities is surface densification. Nonetheless, the microstructural features of the CS + LSP specimens were refined. It is commonly known that when a microstructure is refined, the hardness is increased due to the well-known Hall–Petch relationship, with then can improve the adhesion resistance of a material. To further support this notion, Archard’s theoretical equation for wear [92] suggests that harder surfaces result in less surface deformation under tribo-loading. However, it should be mentioned that the tribological characteristics of materials such as polymers do not necessarily follow this rule, as increases in hardness can promote a brittle-like behavior, thus resulting in greater surface deformation [9395]. Although the brittle-like behavior is not the case for all polymers (depending on their composition), this variation can exist. Nonetheless, when these instances occur for metals (such as in this work), it is expected that there would be a decrease in friction due to the lesser degree of shearing from the tribo-loaded surface. To further explicate, the aforementioned notion essentially suggests that flow stress during tribo-loading is effectively reduced due larger fraction of atoms along the grain boundaries [96]. This explanation further is validated by observing both the change in wear track width and adhesive wear along the wear track surface. Aside from the COF being decreased, the lesser degree of plastic deformation also suggests that fewer asperities are interlocked, thus resulting in a lesser area of contact area and less material transfer [97], as seen in this work.

Lastly, the change in surface texture can also alter the tribological behavior of the pre- and-post-LSP specimens. Generally speaking, the increase in the Sa parameter (due to the asperity roughening effects of LSP [37]) has been reported to be proportional to friction due to the increased rate of shearing along the contacting surfaces [87,98]. From an asperity perspective, the roughened asperities have a greater chance of interlocking with the counter Al6061 material, thus resulting in a greater frictional force. However, when the asperities were smoother, the frictional force decreases. According to the work from Du et al. [99,100], the frictional response of a material is largely influenced by the actual area of contact during sliding alongside the shear strength of the tribo-loaded material. In the case of the smoothened LSP specimens, there were fewer contacting asperities (i.e., lesser contact area) from the counter pin. Additionally, due to the strengthening of their surfaces, the restructured asperities also exhibited a greater shear strength. This was quite evident with the obtained COF plots (Fig. 14), as the frictional response of the tribo-loaded region decreased as the LSP surface became less rough. Especially from a material-transferring perspective, the lessened surface roughness would insinuate that there is a lesser degree of shearing from the contacting asperities, thus resulting in less Al being transferred to the CS surface. With the same idea, it further supports the notion that the greater roughness of CS + I1 resulted in greater Al transfer relative to the other specimens. However, the complexity of the surface’s tribological characteristics (i.e., transfer layer content and frictional coefficient) pre- and-post-peening cannot be simply explained by the aforementioned parameters individually. Rather, more complex factors that relate to all of the aforementioned reasons need to be taken into consideration. In an attempt to relate the change in tribological performance, a new hybrid surface parameter is introduced. Termed as ζF, this parameter serves to correlate the coefficient of friction to amplitude roughness parameters and key surface properties. This parameter is described as
ζFHSPSa(kgfmm2nm)
(5)
where Sa is the arithmetic average surface roughness, H is the surface hardness, and SP is the surface porosity. Figure 17 depicts the plotted values of the newly formed surface parameter to the coefficient of friction. Here, having a higher ζF value is more desirable for friction reduction. Upon closer inspection, it can be seen that when a material exhibits a higher hardness or smaller porous content, there is greater ζF value (i.e., greater frictional resistance), as per the aforementioned discussion. These values have the greatest influence on this coefficient. In the same way, if the surface roughness is lessened, ζF is effectively increased.
Fig. 17
Correlation between the coefficient of friction and the newly established roughness parameters ζF
Fig. 17
Correlation between the coefficient of friction and the newly established roughness parameters ζF
Close modal

The overall mechanisms proposed in this work as schematically shown in Fig. 18. As such, this model suggests three key findings. First, LSP is indeed a viable post-processing technique to augment the surface topography of CS specimens. Second, the influence of LSP can indeed reduce the surface frictional response by pore closure, surface smoothing, and surface hardening. A hybrid surface parameter, ζF, was introduced to relate this relationship. Third, the transfer layer from tribo-loading can also be controlled through the aforementioned conclusions. Collectively, this work acts as a novel foundation for understanding the frictional response and transfer layer mechanisms of post-processed CS coatings, a necessity that is needed to improve their operational performance.

Fig. 18
Conceptualization of the tribological mechanisms of CS 316L SS subjected to LSP
Fig. 18
Conceptualization of the tribological mechanisms of CS 316L SS subjected to LSP
Close modal

5 Conclusions

In this work, the effect of LSP on the surface morphology, microstructure, and tribological performance of CS 316L stainless steel was investigated. The primary findings in this work are as followed:

  • As the laser intensity increased, the degree of surface roughness gradually increased due to the plastic deformation from the plasma-induced shockwaves of the laser. However, after utilizing a 2.54 GW cm−2 laser intensity, the surface roughness decreased. This finding was likely attributed to the high amplitude of compressive stresses along the surface which can result in pore healing.

  • Due to the synergistic combination of severe plastic deformation and rapid heating/cooling, homogeneous nanoscale cellular grains were fabricated along the surface. This was especially evident in higher laser processing intensities.

  • LSP resulted in a refinement in crystallinity, a net increase in dislocation density, and a net increase in strain due to the influence of severe compressive stresses. However, the crystallinity appeared to have an increasing trend whereas the dislocation density had a decreasing trend. Similarly, the strain initially decreased but began increasing as laser intensity became more intense. These findings are attributed to the structural reformation dominant mechanisms of LSP without a coating, as lower intensities were more deformation-based whereas higher intensities were diffusion-based.

  • The COF and degree of Al transfer layer to the CS + LSP specimens were observed to drastically decrease. Findings indicate that this was largely due to the densification of the surface, having desirable surface roughness characteristics (i.e., low surface roughness), and grain refinement due to the LSP process. In an attempt to simplify these key relationships, a new hybrid surface parameter, ζF, was introduced to relate these factors to frictional performance.

Based on the findings of this work, it appears that LSP not only acts as a technique that improves the tribological performance of CS deposits but is also a useful technique to densify the surface. With this work being among the first that studies the effect of LSP on the surface morphology and tribological performance of CS substrates, the authors suggest further investigation into this topic should be conducted. By doing so, many industries which utilize CS as a protective coating will greatly benefit.

Acknowledgment

The authors would like to thank the National Science Foundation (Grant No. CHE-1429768) for allowing the use of the powder X-ray diffractometer.

Conflict of Interest

There are no conflicts of interest.

Data Availability Statement

The data sets generated and supporting the findings of this article are obtainable from the corresponding author upon reasonable request.

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