Surface Hardening of Stainless Steel

Abstract

The addition of nitrogen to stainless steel improves mechanical and corrosion properties. Nitrogen-bearing stainless steel (HNSS) is a new corrosion-resistant alloy class exhibiting better tribological properties. High-pressure and powder metallurgy techniques were developed for the fabrication of HNSS. Solid-state routes allow nitrogen introduction through thermochemical, implantation, or plasma surface treatments. High-temperature gas nitriding (HTGN), carried out in an N2 atmosphere in the 1000°C range, allows N uptake, obtaining thick, ~0.5–1.0 wt.% N austenitic cases. HTGN is different from conventional nitriding, performed in the 500°C range, where intense CrxNy precipitation occurs, impairing the corrosion resistance. Low-temperature plasma nitriding (LTPN) introduces more N in solution, and colossal supersaturated expanded phases (~45 at.%N) are formed. N supersaturation and compressive stresses increase the hardness of the surface layer to 10–14 GPa. Ferritic, martensitic, duplex, and precipitation-hardened stainless steels can be surface-treated by LTPN, obtaining expanded ferrite and martensite. However, single LTPN stainless steel may prematurely fail when submitted to high loading, as the thin and hard expanded layers collapse due to lack of load-bearing capacity. Duplex-nitriding treatment (HTGN + LTPN) results in a thick nitrogen-rich hardened austenite substrate layer, granting mechanical support and adhesion to the expanded austenite layer.

1. Introduction

Since the beginning of the twentieth century, stainless steel has been developed to improve the corrosion resistance of parts in contact with corrosive and oxidative media. These corrosion-resistant alloys have been used in the chemical, petrochemical, automotive, aeronautical, food, medical, and construction industries. Chromium, above 11 wt.%, grants corrosion resistance by forming a nanometric thin and adherent Cr₂O₃ passive layer. When exposed to oxygen, whether in the air or water, this layer prevents corrosion by isolating the alloy from contact with the oxidizing media.

However, chlorine and chlorine ions may damage the passive layer favoring stainless steel’s crevice, pitting, and stress corrosion cracking. Mo additions are very effective in improving the resistance to damage of the passive layer by chlorine, although it negatively influences the final price of the stainless steel. On the other hand, N has been thoroughly investigated, since the 1980s, as an alloying element with great potential for protecting the passive layer against damage, being abundant in nature (21 wt.% in the atmosphere), and giving a cost-effective solution for surface protection against corrosion.

Controlled addition of nitrogen to stainless steel has been encouraged over the last three decades due to the possibility of improving the surface properties (not only the corrosion but also the tribological and mechanical properties). High-pressure and powder metallurgy techniques were developed for medium and large-scale fabrication of high nitrogen steels (HNS). Still, in general, these procedures are costly and require sophisticated equipment. Nitrogen-bearing stainless steel is a new class of corrosion-resistant alloys, exhibiting much better surface properties, better corrosion, and wear resistance associated with good bulk mechanical properties: very high strength, good ductility, and toughness. Therefore, considerable emphasis has been placed on liquid and solid-state routes to produce high-performance, low-cost nitrogen-alloyed stainless steels. The liquid state processing routes demand high-pressure metallurgy, which is laborious, demands special equipment, and is costly. In the solid-state production routes, the steel surface and near-surface regions are nitrogen alloyed through thermochemical, implantation, plasma, or laser techniques.

Diffusion surface treatments have been extensively studied and have become, for many applications, current industrial practice. The diffusion of nitrogen and carbon toward the core increases the surface hardness and wear resistance. However, nitrogen and carbon must remain in solid solution. Precipitation of chromium-rich carbides or nitrides reduces the chromium content in the metal matrix, preventing the formation of a continuous passive layer and harming the corrosion resistance of the steel.

2. High-temperature diffusion surface treatment

Berns [1] proposed, in the first half of the 1990s, carrying out a high-temperature nitriding process by exposing austenite to an N₂ atmosphere. Nitrogen dissolves in austenite up to the solubility limit during the high-temperature nitrogen treatment. Nitrogen solubility in austenite is much greater than in the BCC phases. Then, by maintaining stainless steel in a furnace containing a pure N₂ gas atmosphere, the nitrogen equilibrium between the furnace atmosphere and the alloy can be attained. According to Sieverts’ law [2], nitrogen can reach contents up to 1 wt.% in solution. Thermocalc [3] phase diagrams considering the N₂ gas phase were calculated to predict the N₂ content in equilibrium in austenite as a function of temperature and partial pressure, as shown in Figure 1.

Figure 1.

TPT diagram relating nitriding parameters (nitrogen temperature and partial pressure) with microstructure, nitrogen content, and martensitic layer depth for 3 h gas nitriding treatments at high temperature for an AISI 410S steel [4].

The high-temperature nitriding treatment consists of a case hardening that enriches the stainless steel’s surface with nitrogen contents up to 1 wt.%, to a depth of 1–2 mm. Berns [5] named this process solution nitriding (SN). After this pioneer proposal, several research works have followed on studying this solid-state route for introducing high N₂ contents in solution in austenite, being called high-temperature-gas-nitriding—HTGN [6] or high-temperature-solution-nitriding HTSN [7].

The amount of nitrogen dissolved in austenite, in equilibrium with pure N₂ gas atmosphere, increases with decreasing temperature and pressure, as shown in Figure 2 [6].

Figure 2.

Fe—13% Cr—N isopleths (a) not considering the gas phase as an equilibrium one, with N₂ isobars overlaid and (b) considering the N₂ gas phase as an equilibrium one [6].

Berns [5] envisaged different possibilities of obtaining tailored engineered stainless steels depending on the composition and surface treatment. Therefore, austenitic stainless steels can be HTGN, obtaining a fully austenitic case with an outermost N content of 0.48 wt.%N and case depths of up to 1 mm, with a hardness variation from 1.95 GPa in the low N core to 3.17 GPa in the 0.48 wt.%N case [8]. Martensitic stainless steels can be HTGN, obtaining a much harder 0.4 wt.%N martensitic cases 725 HV hard [9]. Extra-low carbon (0.017 wt.%C) dual-phase stainless steel (α + Martensite) may form a fully martensitic case 550 HV hard, after HTGN [10]. Finally, an UNS 31803 ferritic/austenitic duplex stainless steel can be hardened by HTGN, achieving a fully austenitic layer near the surface due to enrichment in austenite stabilizer element (N), as shown in Figure 3 [11]. Excess of diffused nitrogen causes a solid solution hardening effect, proportional to its content, reaching a maximum value of 330 HV at maximum concentration, as shown in Figure 4 [12]. It is worth noting that the N absorption and diffusion on the surface during the HTGN process induce phase transformations, resulting in microstructural gradients from the surface to the core and corresponding microhardness gradients.

Figure 3.

UNS S31803 duplex stainless steel HTGN at 1200°C [11].

Figure 4.

Microhardness gradient from the low nitrogen duplex ferritic-austenitic core toward the fully austenitic 0.8 wt.%N surface [12].

Tschiptschin [13], using this concept, proposed a Powder Metallurgy route to enrich a ferritic stainless steel powder (0.02 wt.%C, 16.2 wt.%Cr, and 0.81 wt.%Mo), exposing the powder particles at high temperatures (1100°C and 1200°C) to N₂ gas atmosphere. The N enriched austenitic powder transforms during quenching to martensite, becoming very hard. One of the main challenges in this HNS production route is obtaining fully dense components with uniform nitrogen content in volume and excellent surface properties. A uniform nitrogen distribution leads to a more homogeneous microstructure and better mechanical properties. Figure 5 shows the amount of nitrogen as a function of temperature. According to Sieverts’ law [2], increasing temperature decreases the amount of nitrogen content of the obtained alloy. High temperatures are necessary to grant that all the nitrogen is dissolved in austenite, avoiding the precipitation of chromium nitrides.

Figure 5.

Nitrogen content as a function of temperature and N₂ pressure for an AISI 434L ferritic stainless steel [13].

In this route, high-nitrogen (0.66 wt.%N) martensitic stainless steel could be obtained by high-temperature gas nitriding an AISI 434L ferritic stainless steel powder, compressing the high nitrogen powder to near net shape parts, followed by hot isostatic pressing and proceeding with a 1200°C quenching and a 200°C tempering treatment. As a result, the obtained hipped material showed high hardness and much better corrosion resistance, measured in potentiodynamic polarization tests carried out in 0.5 M H₂SO₄ + 3.5% NaCl, as shown in Figure 6.

Figure 6.

Cyclic polarization curves for a 0.66 wt.%N martensitic stainless steel in different stages of fabrication. Solution 0.5 M H₂SO₄ + 3.5% NaCl. S/N: Sintered/nitrided, HIP: Hot isostatic pressed, HT: Heat treated (quenched and tempered at 200°C) [13].

3. Low-temperature diffusion surface treatment

Conventional gaseous or liquid nitriding processes are traditionally carried out at temperatures above 520°C. However, this process temperature is a limiting factor, considering that when the nitriding of stainless steels is conducted above 500°C, intense precipitation of chromium nitrides and carbides occurs in the diffusion zone, which, despite substantially increasing the hardness, greatly compromises corrosion resistance [14, 15, 16, 17, 18].

The diffusion temperature is the main control parameter to prevent chromium nitrides and chromium carbides precipitation. Precipitation of chromium carbides and nitrides requires substitutional diffusion, which only occurs at temperatures higher than 500°C. Zhang and Bell [14] and Ichii et al. [19] pioneered the study and development of stainless steel’s nitrogen and carbon diffusion processes in low temperatures. The process temperature must be selected, not too low, to allow intense diffusion of the C and N interstitial elements but not high enough to permit substitutional diffusion. At these low temperatures, the chromium substitutional element’s mobility is sufficiently reduced to inhibit the nucleation and growth of nitrides and/or carbides. Under these conditions, the matrix becomes continuously and increasingly enriched by the interstitial element, promoting a non-equilibrium saturation of the crystalline lattice and stabilizing expanded phases formed in the diffusion layer.

Interstitial supersaturation in the diffusion zone contributes to: (i) formation of interstitially supersaturated phases, (ii) intense interstitial hardening, as a consequence of the colossal amount of interstitial element and its stress fields, contributing to increasing the wear resistance without compromising corrosion resistance [20, 21, 22, 23, 24], and (iii) generation of residual compressive stresses in the expanded layer as a result of the restrictive effect of the diffusion-free substrate, which improves fatigue properties [24, 25].

When performing the X-Ray Diffraction of these supersaturated layers, it is observed that the matrix peaks are shifted to lower 2θ angles and show a greater FWHM - full-width at half-maximum height than the peaks of the unenriched matrix phase. This displacement and broadening of the peaks indicate elastic deformation due to the expansion of the crystalline lattice. Zhang and Bell [14] named this phase “Expanded Phase” due to the expansion of the lattice parameters of the crystalline unit cell. On the other hand, Ichii et al. [19] called this phase S-Phase due to the “shifting” to lower angles of the XRD peaks.

Bell and Chen [26] and Sun [27] presented a limit curve for precipitation of chromium nitrides and carbides as a function of temperature and time of plasma diffusion process for an austenitic AISI 316 L stainless steel. “Nitrogen Expanded Austenite (γ_N)” and “Carbon Expanded Austenite (γ_C)” are formed during nitriding or carburizing for temperatures and times below the limit curves shown in Figure 7.

Figure 7.

Limit curves for precipitation of chromium nitrides or chromium carbides in austenite as a function of temperature and time of plasma diffusion process [26, 27].

Expanded phases obtained at diffusion temperatures between 350°C and 430°C are responsible for surface hardening [20, 21, 28, 29, 30, 31]. This hardening can be obtained in all stainless steel families with the formation of different phases expanded by nitrogen and/or carbon [20, 32, 33, 34, 35, 36]. Table 1 shows the different expanded phases formed by low-temperature plasma diffusion surface treatment and their hardening characteristics for different families of stainless steels.

A passivating Cr₂O₃ film formed on the surface of the parts to be nitrided prevents nitrogen or carbon from entering stainless steel. Thus, the passive film’s mechanical or chemical removal process should be employed before diffusion. The chemical removal of the passive film by acid pickling may compromise the surface finish of the parts or maybe potential damage to the operators’ health or equipment. Currently, modern low-temperature gas nitriding processes still use acid pickling [37] for depassivation of the Cr₂O₃ layer during exposure of the parts’ surface to atmospheres containing halides (NF₃ or HCl) has been carried out in Low-temperature gas carburizing [38]. Activation of the parts’ surface by nickel plating to prevent repassivation by catalytic decomposition of NH₃ gas [39] has also been used.

Activation of the surface by “sputtering” in H₂, under high voltage and low pressure—[40], use not only the kinetic energy of the ions but also the reducing character of hydrogen [41], preserving the surface quality of the parts being nitrided.

3.1 Austenitic stainless steels

The behavior of nitriding at high temperatures, above 500°C, and at low temperatures, below 420°C, mainly affects the corrosion resistance of the nitrided surface [14]. Austenitic stainless steels cannot be nitrided conventionally at temperatures close to 500–550°C due to intense precipitation of CrN and Cr₂N chromium nitrides in the diffusion zone [42, 43, 44, 45]. The precipitation of these nitrides increases the surface hardness but greatly decreases the corrosion resistance due to chromium removal from the solid solution in the matrix. Nitriding must be carried out below 430°C in order to avoid precipitation of nitrides. In this low-temperature nitriding process, generally between 380°C and 420°C, the diffusion kinetics of the chromium substitutional element is significantly reduced, which inhibits the formation of chromium nitrides. The increasing diffusion of nitrogen in the austenite generates a supersaturated solid solution that expands the CFC crystalline lattice and forms the metastable phase called Expanded Austenite—(γ_N) [46, 47, 48]. The formation of expanded austenite promotes an increase in surface hardness without compromising corrosion resistance [20, 49, 50].

Figure 8 shows the microstructures obtained by Bruno et al. [50] for AISI 316 L steel after nitriding at temperatures of 550°C (a) and 380°C (b). When this steel is nitrided at 550°C, the nitrided surface becomes dark and severely etched, which denotes the loss of corrosion resistance in this region under the action of Villela’s reagent. When nitriding is carried out at 380°C, the nitrided surface appears as a white layer, and the non-nitrided matrix shows a microstructure very similar to a typical austenitic steel microstructure. In this low-temperature nitriding condition, the Marble metallographic reagent does not etch the nitride layer, only the matrix, which indicates a better corrosion resistance of the nitrided surface. This corrosion resistance behavior against metallographic reagents results from the formation mechanism of expanded phases on the nitrided surface. The X-ray diffractograms in Figure 9 show the non-nitrided condition after nitriding at 380°C (a) and 550°C (b). In the non-nitrided condition, only FCC austenite peaks are present. When nitriding at 380°C, FCC austenite peaks are shifted to the left and become broader, resulting from residual compression stresses and distortion of the crystalline lattice caused by nitrogen supersaturation. When expanded austenite is formed on the surface of the nitride specimen, the corrosion resistance is maintained or even improved compared to the non-nitrided specimen. On the other hand, when nitriding is carried out at 550°C, several CrN and Cr₂N diffraction peaks show up. Chromium nitride precipitation induces depletion of the Cr content of the metallic matrix and is ​​responsible for the loss of corrosion resistance of the nitrided surface.

Figure 8.

Images of the nitrided surfaces after nitriding at (a) 380°C and (b) 550°C. Bruno et al. [50].

Figure 9.

XRD spectra for AISI 316L steel before and after nitriding [50].

The expansion of the FCC crystal lattice and the increase of the lattice parameter, which occurs when expanded austenite is formed, are shown in Figure 9. Expanded austenite peaks are shifted to the left, and the volume variation is close to 10% [20]. Strain-free FCC austenite has a lattice parameter equal to 0.359 nm (ICDD^® Card 00–033-0397). After plasma nitriding, the lattice parameter in expanded austenite increases to 0.375 nm, corresponding to a calculated nitrogen content at a supersaturation equal to 34.6% atomic or approximately 8.5% by mass. These estimations do not consider the contribution of the residual stresses in shifting the diffraction peaks to the left) [51, 52]. Expanded austenite is responsible for the increased surface hardness up to 7 times over the original hardness, as shown in Figure 10 [20].

Figure 10.

Surface hardening due to the formation of expanded austenite after plasma nitriding AISI 316L stainless steel at 400°C [20].

The interstitial supersaturation of the matrix may be due to nitrogen diffusion in nitriding or carbon diffusion [25, 26, 27, 28, 29, 30, 31, 32, 33, 34, 35, 36, 37, 38, 39, 40, 41, 42, 43, 44, 45, 46, 47, 48, 49, 50, 51, 52, 53, 54] upon plasma carburizing. The surface treatment may comprise both nitrogen and carbon diffusion, and the plasma treatment is called nitrocarburizing or just carbon for plasma carburizing. These treatments may be carried out at low temperatures, below 430°C for nitrocarburizing and below 500°C for carburizing, avoiding carbide or nitride precipitation. Figure 11 shows the microstructures after (a) nitriding, consisting of a monolayer of austenite expanded by nitrogen (γ_N); (b) nitrocarburization consisting of a double layer composed of austenite expanded by nitrogen (γ_N) in the outer region and carbon expanded austenite (γ_C) between the first layer and the matrix; and (c) carburizing consisting of a carbon-expanded monolayer of austenite (γ_C).

Figure 11.

Microstructures of austenitic stainless steel after (a) nitriding, (b) nitrocarburizing, and (c) carburizing at 400°C. scanning electron microscopy [54].

In nitrocarburizing and carburizing, colossal interstitial supersaturation leads to expansion of the crystalline lattice, generating the expanded phases “γ_N” and “γ_C.” Table 2 shows the expansion characteristics of the austenite FCC crystalline lattice under each condition and the dissolved nitrogen content in the supersaturated condition. The volume expansion of the FCC lattice, indicated by the ratio ∆a/a, is responsible for the hardening and the generation of residual compressive stresses.

Steel	Classification	Expanded phase	Symbol	Typical hardness (HV)
Austenitic	AISI 316	Expanded Austenite	γN	1400
Martensitic	AISI 420	Expanded Martensite	α′N	1000
Precipitation-hardening	17-4PH	Expanded Martensite	α′N	1000
Ferritic	AISI 410S	Expanded Ferrite	αN	1200
Duplex	AISI F51	Expanded Ferrite	αN	1000
		Expanded Austenite	γN	1000

Table 1.

Expanded phases formed during low-temperature nitriding of stainless steels.

	Phase	a (nm)	∆a/a (%)	Concentration, N-C (calculated)
Non-nitrided	γ	0.3589	—	—
Nitrided (400°C)	γN	0.3924	9.32	37.2 at.%N
Nitrocarburized	γN	0.3916	9.11	36.3 at.%N
(400°C)	γC	N.D.	N.D.	N.D.
Carburized (400°C)	γC	0.3627	1.05	6.2 at.%C
Carburized (480°C)	γC	0.3684	2.64	13.9 at.%C

Table 2.

Lattice parameters, lattice expansion, and calculated interstitial content at supersaturation for the expanded austenite layers in AISI 316L steel.

Figure 12 shows the relationship between carbon supersaturation in the FCC crystalline lattice and residual stresses in the case-hardened surface of AISI 316 L austenitic stainless steel. Both parameters gradually decrease toward the nucleus [53] due to supersaturation and the generation of residual compressive stresses on the treated surfaces. Figure 13 shows the high potential for surface hardening for the three types of treatment [54].

Figure 12.

Carbon pickup and compressive residual stresses on the surface of AISI 316L stainless steel after low temperature carburizing to colossal carbon enrichment [53].

Figure 13.

Maximum hardness on AISI 316L austenitic stainless steel surface upon plasma nitriding, nitrocarburizing, and carburizing at 400°C [54].

3.2 Martensitic stainless steels

Martensitic stainless steels behave similarly to austenitic stainless steels concerning the formation mechanisms of the nitrided surface at different process temperatures. Figure 14 shows the microstructures of AISI 420 steel after plasma nitriding at 380°C and 550°C. At 550°C, the diffusion zone is heavily darkened due to a severe etching by Villela’s reagent, but when nitriding is carried out at 380°C, the diffusion zone practically remains unchanged compared to the tempered martensite matrix. The darkening of the nitrided layer denotes loss of corrosion resistance due to the nitriding process, while the unetched nitrided layer indicates that the corrosion resistance is maintained in low-temperature nitriding [33].

Figure 14.

Microstructures of nitrided surfaces at (a) 550°C and (b) 380°C [33].

In Figure 15, the X-ray diffraction maps show AISI 420 steel before and after plasma nitriding [33]. After quenching and tempering, only the peaks referring to the tempered martensite (α′) are observed. When nitriding is carried out at 380°C, the tempered martensite peaks give way to the expanded tempered martensite peak (α′_N). Peaks corresponding to iron nitrides, Fe₃N, and Fe₄N are also observed. This result shows that this temperature is low enough to inhibit the diffusion of chromium, preventing the precipitation of CrN and Cr₂N nitrides. Avoiding the precipitation of chromium nitrides at low temperatures is responsible for maintaining corrosion resistance. The CrN and Cr₂N chromium nitrides diffraction peaks that appear after nitriding at 550°C show intense precipitation of chromium compound and chromium depletion of the matrix, responsible for the decrease in the corrosion resistance of the nitrided surface.

Figure 15.

XRD spectra for AISI 420 steel before and after nitriding [33].

Figure 16 shows the corrosion rate of a 380°C nitrided AISI 420 steel specimen when subjected to an immersion test in an aqueous solution with 10% HCl for 120 h. In the quenched and tempered condition, the corrosion rate after nitriding is lower than the non-nitrided material due to the higher nitrogen concentration on the nitrided surface.

Figure 16.

The corrosion rate of AISI 420 steel in aqueous solution with 10% HCl for 120 h before and after plasma nitriding at 380°C for 20 h. Author: Unpublished.

While hardening occurs due to the precipitation of chromium nitrides, in the 550°C plasma nitriding treatment and in the 380°C nitriding treatment, the nitrided surface hardens due to the formation of expanded tempered martensite (γ′N), which induces compressive residual stresses.

Figure 17 shows that, compared to the quenched and tempered matrix, with 590 HV, the low-temperature nitriding plasma treatment (380°C) promotes hardening near 1000 HV. For the 550°C nitriding, the hardening nearly reaches 1300 HV. Despite the lower hardening in the nitriding treatment at 380°C, this condition should be preferentially used, as it combines hardening and good corrosion resistance.

Figure 17.

Maximum hardness after plasma nitriding of AISI 420 steel [32].

Another important factor related to the hardening characteristic is the transverse hardness profile obtained in these two conditions, Figure 18. For the 550°C nitriding treatment, the transverse hardening profile shows a maximum hardness level throughout the diffusion zone with an abrupt drop at the matrix interface [32, 55, 56]. A very steep hardness gradient is not appropriate to withstand mechanical shear stresses found during sliding. Furthermore, exposing the steel to high nitriding temperatures causes a decrease in core hardness by an over-tempering effect [32]. When low-temperature nitriding is carried out, despite the lower maximum hardness, the transverse hardening profile is diffuse, with no decrease in core hardness, and suitable for most different applications.

Figure 18.

Transverse hardening profiles after plasma nitriding of AISI 420 steel at 380°C and 550°C [32].

The surface hardening promoted in the low-temperature plasma nitriding treatment is responsible for increasing the tribological properties [57]. Figure 19 compares the scratch resistance of an AISI 410 martensitic stainless steel: (a) non-nitrided, quenched, and tempered to a 40 HRC hardness; (b) plasma nitrided at 400°C. The scratch path in the non-nitrided condition is thicker and more profound than in the nitrided condition and presents deformation in its surroundings. Table 3 shows that the scratch severity is at least half of the non-nitrided condition for the scratch track’s depth and thickness in the nitrided condition.

Figure 19.

Scratches made under constant load on the surface of an AISI 410 steel in the (a) quenched and tempered and (b) after plasma nitriding at 400°C conditions [57].

Condition	Hardness (HV0.01)	Width (μm)	Depth (μm)
Non-nitrided	388	90	26
Nitrided	1275	43	12

Table 3.

Scratch width and depth for non-nitrided and 400°C plasma nitrided AISI 410 stainless steel.

Figure 20 compares the cavitation resistance of non-nitrided and 400° plasma nitrided AISI 410 stainless steel in a test [58]. One can see that the mass loss of the low-temperature plasma nitrided specimens lost 40 times less mass than the non-nitrided specimen.

Figure 20.

Mass loss during cavitation tests of an AISI 410 steel in the quenched and tempered and 400°C plasma nitrided conditions [58].

Martensitic stainless steels can also be nitrocarburized or carburized [59, 60, 61, 62]. Nitrocarburization of 420 martensitic stainless steel carried out at 450°C for 4 h can achieve a surface hardening close to 1280 HV with a layer composed of nitrogen and carbon expanded martensite (γ′_NC) and Fe₃C/Fe_2–3(CN) type precipitates. Nitriding at lower temperatures avoids these precipitates in the layer. Figure 21 shows the hardness profile of the martensitic stainless steel after plasma hardening at 450°C for 4 h, with a maximum hardening potential of 800 HV and a hardening depth in the diffusion zone close to 0.040 mm [61].

Figure 21.

Hardness profile for a 450°C (4 h) plasma nitrided AISI 420 martensitic stainless steel [61].

3.3 Precipitation hardening stainless steels

The plasma nitriding process for PH precipitation-hardening stainless steels should preferably be carried out at temperatures equal to or below the aging temperature of the steel part. PH steels are aged at different temperatures, specified according to the final desired mechanical properties. A proper selection of the nitriding temperature allows for reaching the desired surface hardness without compromising the quenched pus tempered hardness achieved during aging. The plasma nitriding treatment of PH steel components may be carried out at lower or higher temperatures depending on the application and the operating conditions.

Figure 22(a) shows the nitrided layer and the resulting hardening of aged 17–4PH steel after 4 h at 550°C plasma nitriding treatment. For this condition, the nitrided layer is composed of a diffusion zone formed by the precipitation of iron and chromium nitrides. The precipitation of these nitrides promotes an intense surface hardening, capable of raising the surface hardness to values close to 1300 Vickers [34]. Figure 22(b) shows that the 35.3 HRC hardness of the substrate, previously aged for 4 h at 552°C, condition H1025 (AMS 5643 2013), practically remained unchanged after nitriding at 550°C/4 h with a measured value of 34.4 HRC [34]. Figure 23 shows a gentle hardness profile and a nitriding depth close to 0.65 mm.

Figure 22.

Nitrided layer (a) and hardening characteristic (b) of 17–4PH steel after plasma nitriding in DC-plasma at 550°C [34].

Figure 23.

Transverse hardening profile of 17–4PH steel after DC plasma nitriding at 550°C [34].

Nitriding of PH steels can also be performed at lower temperatures so as not to affect corrosion resistance [63]. Figure 24 shows the nitrided layer of 17–4PH steel after plasma nitriding at 400°C using the active screen technique. It is observed that the nitrided layer is white and not etched by Villela’s reagent, unlike the nitrided layer at 550°C, which is dark and severely etched by Villela’s reagent. This difference in behavior is related to the nitriding mechanisms. As hardening in nitriding at 550°C occurs with the precipitation of chromium nitrides, corrosion resistance decreases as the matrix is depleted in chromium. When plasma nitriding is carried out at 400°C, hardening occurs by forming a nitrogen supersaturated layer of expanded martensite (α′_N) without nitrides precipitation, reaching 1130 Vickers.

Figure 24.

Hardness variation after plasma nitriding of 17–4PH stainless steel α′_N expanded martensite layer and Vickers hardness before and after ASPN at 400°C [63].

3.4 Ferritic and duplex stainless steels

Low-temperature plasma nitriding of ferritic and duplex stainless steels is still being developed and is not commercially available yet. However, many reports promise good results for use in most different components and applications. When ferritic stainless steels are nitrided at low temperatures, precipitation of chromium nitrides is avoided [35, 64]. Figure 25 shows the microstructure of a plasma nitrided AISI 410S stainless steel with a layer of expanded ferrite (α_N) containing Fe₃N iron nitrides. Shifted to the left, expanded ferrite (α_N) peaks appear on the X-ray diffraction pattern of the nitrided layer, Figure 26. Vertical dashed lines indicate the positions of ferrite peaks in the matrix. Besides, Fe₃N peaks were also detected. The absence of chromium nitrides in the nitrided layer grants the corrosion resistance of the nitrided surface. Figure 27 shows the hardening obtained in the nitriding by comparing the maximum surface hardness obtained and the transversal hardening profile of the nitrided surface.

Figure 25.

Microstructure of the nitrided surface after plasma nitriding AISI 410S stainless steel at 400°C [64].

Figure 26.

XR diffraction pattern of the surface after plasma nitriding AISI 410S stainless steel at 400°C [64].

Figure 27.

Surface hardness and hardness profile of an AISI 410S stainless steel after plasma nitriding at 400°C [64].

Corrosion resistance testing carried out by immersion in 3% FeCl₃ aqueous solution, for 88 h, at room temperature showed a better performance of the nitrided specimens concerning the non-nitrided ones, Figure 28. When the steel is nitrided at a low temperature (N400°C), the corrosion properties are not changed compared to the non-nitrided condition. However, high-temperature nitriding (N530°C) promotes a significant loss of corrosion resistance compared to the other two conditions [65].

Figure 28.

Mass loss and corrosion rates of non-nitrided and plasma nitrided AISI 410S stainless steel during immersion in 3% FeCl₃ aqueous solution for 88 h at room temperature [65].

Duplex stainless steels’ microstructure is composed of austenite and ferrite in approximately equal proportions. In this condition, low-temperature nitriding leads to the formation of expanded austenite (γ_N) and expanded ferrite (α_N) on top of ferrite and austenite strings, respectively [36, 66, 67]. Figure 29 shows the microstructure on the surface of type 2205 duplex stainless steel after plasma nitriding at 400°C. The austenite and ferrite bands and the formation of the respective expanded phases on the nitrided surface are observed in the photomicrograph [67]. The X-ray diffraction pattern in Figure 30 shows the initial phases’ peaks and shifted to the left, the respective peaks of the nitrogen-expanded phases. Fe₃N iron nitrides were also detected [67]. Consequently, the formation of expanded austenite and expanded ferrite on the surface led to an intense hardening of the nitrided surface, as shown in Figure 31.

Figure 29.

Microstructure of 2205 duplex stainless steel, after low-temperature plasma nitriding. Expanded ferrite and expanded austenite [67].

Figure 30.

XR-diffraction pattern of 400°C plasma nitrided 2205 duplex stainless steel [67].

Figure 31.

Hardness of ferrite and austenite and expanded ferrite and expanded austenite in 400°C plasma nitrided 2205 duplex stainless steel [36].

4. Duplex diffusion surface treatments

Duplex stainless steels can also be subjected to a duplex nitriding treatment consisting of a combination of high-temperature and low-temperature diffusion treatments. Firstly, an HTGN—high-temperature gas nitriding is carried out at 1200°C, followed by a 400°C LTPN—low-temperature plasma nitriding aiming for a better load-bearing capacity. In the high-temperature nitriding treatment, nitrogen is introduced on the surface of the steel, shifting the phase equilibrium so that ferrite stringers are transformed to austenite, thus forming a 100 μm thick fully γ layer, raising the hardness from 280 to 330 HV. Subsequently, the LTPN—low-temperature plasma nitriding diffusion treatment leads to a continuous and homogeneous layer of expanded austenite (γ_N), 1200 HV hard, on top of the austenite layer [11]. Figure 32 shows the microstructure of the duplex nitrided 2205 steel. This microstructure grants greater load-bearing capacity than the single plasma nitriding treatment, and the alloy’s performance under cavitation-erosion is much better, as shown in Figure 33.

Figure 32.

Microstructure of duplex treated 2205 duplex stainless steel (HTGN + LTPN) [11].

Figure 33.

Mass loss during a cavitation-erosion test. 2205 duplex stainless steel non-nitrided, HTGN, and duplex nitrided (HTGN + LTPN) [11].

Duplex treatments (HTGN + LTPN) can considerably increase the tribological properties of the surface. Figure 33 shows the mass loss results during a 2205 duplex-stainless-steel cavitation-erosion test. The mass loss decreases after (HTGN) high-temperature gas nitriding, forming a thick high-nitrogen austenite layer. In the duplex treatment, the mass loss is almost null for testing times up to 64 h due to forming a 1300 HV hard expanded austenite layer sustained by a harder substrate [11].

5. Plasma diffusion surface treatment

A plasma diffusion technology, which combines a plasma nitriding treatment with a subsequent solubilization treatment, was proposed in pioneering work by Pinedo et al. for solid-state alloying [68, 69]. A 1 mm thick sample of AISI 316 L austenitic stainless steel, plasma nitrided at 470°C for 12 h, in a 1N₂:1H₂ gas mixture, formed a 60 μm deep, 1290 HV hard nitrided layer composed of γ + CrN + Cr₂N. After nitriding, the material was solubilized at 1150°C and cooled in water to promote the diffusion and homogeneous redistribution of nitrogen through the sheet’s cross-section. After solubilizing, a 0.80 wt.%N, homogeneous nitrogen content was found throughout the thickness, consistent with Thermocalc^© predictions. Figure 34 shows the nitrogen enrichment profile on the nitrided surface, reaching a maximum content of 10 wt.%. Figure 35 shows the hardness profile along the sheet thickness showing a homogeneous increase in hardness from 200 to 300 HV, achieving complete hardening of the cross-section.

Figure 34.

Compositional profile of nitrogen obtained by GDOES after plasma nitriding of AISI 316 L steel [69].

Figure 35.

Transverse hardening of AISI 316 steel nitrided under plasma and solubilized at 1150°C [69].

Recently [70] Berton et al. applied the same principle for an AISI 409 ferritic stainless steel. AISI 409 specimens 0.8N₂:0.2H₂ plasma nitrided at 510°C for 2 h and later solubilized at 1100°C for 1 h for nitrogen diffusion in the ferritic matrix. Once the steel was solubilized, it was subjected to quenching and tempering treatments to obtain a high surface hardness tempered martensite layer. Figure 36 presents the nitrogen profile along the cross-section for the solubilized (NS) and tempered conditions from 950°C (Q950) and 1050°C (Q1050), showing a maximum effective enrichment of the order of 1% by mass. This enrichment after the complete quenching and tempering cycles promotes a maximum rise in surface hardness close to 4x the core hardness at transverse hardening depths of up to 600 μm, as shown in Figure 37.

Figure 36.

Compositional profile of nitrogen determined by WDS in the diffused condition (NS) and after quenching from 950°C (NS-Q950) and 1050°C (NSQ1050) [70].

Figure 37.

Transverse hardening profiles of AISI 409 steel under conditions; solubilized (NS) and after quenching at 950°C (Q950) and 1050°C (Q1050), and tempering for 1 h at 250 (T2), 450 (T4), and 650°C (T6). The untreated condition is shown for comparison [70].

6. Conclusion

Although stainless steels are designed to sustain distinctly superior corrosion resistance in a wide range of aggressive environments, these steels do not show enough wear resistance in many highly demanding tribological conditions, harming performance and service life.

Diffusion-based surface hardening processes are an alternative to increasing the wear resistance of stainless steel without compromising its corrosion properties. Accordingly, different thermochemical solutions of diffusion surface treatment of stainless steels were presented; (i) high-temperature gas nitriding, (ii) low-temperature plasma or gas nitriding, nitrocarburizing, or carburizing, (iii) duplex treatment combining high and low-temperature nitriding treatments, and (iv) solid-state annealing, to promote surface hardening and maintain or even raise, the corrosion resistance of these materials.

Surface hardening by diffusion thermochemical processes is an efficient strategy to produce tailor-made surfaces with improved mechanical strength and wear resistance, applicable to all classes of stainless steel.

Acknowledgments

The authors acknowledge the financial support of the Fundação de Amparo à Pesquisa do Estado de São Paulo, process FAPESP 2019/18572-7.