Kinetics of the UNS S32750 Super Duplex Stainless Steel Low-Temperature Plasma Nitriding

Abstract

This work reports an experimental investigation on the thermodynamic and kinetic view-point of the nitrided layer growth behavior, carried out separately at ferrite and austenite phase grains, for the UNS S32750 super duplex stainless steel (SDSS). For this purpose, and motivated by the DSSs distinct behavior showing formation of two- or single-phase nitrided layers, which apparently depend on the steel substrate chemical composition, two series of treatment were studied here, one varying the nitriding temperature on the 300-450 °C range for 4 h time, and other one varying the time on the 2-8 h range at 350 °C. Microstructural characterization of studied layers by means of SEM and EDS analysis, XRD and microhardness measurements showed precipitation-free nitrogen-expanded austenite layer formation on both steel phases for all samples treated up to 400 °C. The nitriding kinetics study showed that the layer thickness on both steel phases is proportional to the square root of the treatment time and follows an Arrhenius law for the studied treatment temperatures. The nitrogen diffusion activation energy, separately determined from the layer growth on each steel phase, was 115±3.2 and 120±5.4 kJmol^-1, for austenite and ferrite grains, respectively.

Keywords:
UNS S32750 super duplex stainless steel; low-temperature nitriding; nitriding kinetics; plasma nitriding; nitrogen-expanded austenite

1. Introduction

Duplex stainless steels (DSSs) have been extensively used in several industrial segments, like shipbuilding, offshore, chemical, paper and pulp, petrochemical, desalination plants and oil and gas, where high corrosion resistance (namely pitting, crevice and stress corrosion resistance) and significant mechanical properties, conferred by its ferritic-austenitic structure, are required¹. DSSs typically contain, in wt. %, 18.5-27 Cr, 4-8 Ni and 2-6 Mo and have, while solution-treated, ferrite-austenite (half to half) microstructure with preferential partitioning of chromium and molybdenum in ferrite (α) grains, and nickel in austenite (γ) grains². One reason for the growing research and development interest in DSSs lies in the fact that they tend to be ideal materials for use in marine engineering equipment³-⁵. Allied to it, DSSs present good potential to replace the costly nickel-based and corrosion-resistant alloys, like austenitic stainless steels (ASSs) that typically contain 8-30 wt.% Ni in its compositions, mainly due to the strong nickel market price volatility¹. This explains why the production of DSSs tends to grow at the expense of a decrease in the production of ASSs⁶. On the other hand, despite of good overall properties, DSSs tends to present low wear resistance, thus weak tribological behavior⁷, which is limiting factor for expanding their use.

In applications where surface properties are important, surface treatments constitute an attractive technological tool to improve the DSSs tribological behavior, increasing the components lifespan⁸. The benefits of the nitrogen-alloying by plasma treatments at the surface of ferritic⁹,¹⁰ as well as austenitic¹¹-¹⁵ stainless steels are well established. In contrast, the response of the DSSs to plasma nitriding was less explored, being far from completely understood, deserving additional research efforts. The existing research indicates that the application of nitriding temperatures above 500 °C, promote the formation of Cr and Fe nitrides, increasing the DSSs hardness and wear resistance, but reducing their corrosion resistance¹⁶. In its turn, the application of low-temperature (paraequilibrium) nitriding, typically performed below 450°C, enables the obtainment of chromium nitride precipitation-free hardened surfaces, presenting a nitrogen-expanded supersaturated structure, which displays better performance to failure by corrosion and wear when compared to non-nitrided material¹⁷-²⁰.

Table 1 presents a compilation of the main papers on the literature about the low-temperature plasma nitriding (LTPN) of DSSs⁷,⁸,¹⁷-²⁸. Regarding the nitrided layer formation mechanisms and its phase constitution, a portion of these research point out that treated layer is constituted of a two-phases γ_N (nitrogen-expanded austenite) and α_N (nitrogen-expanded ferrite) structure⁷,⁸,²¹,²²,²⁷,²⁸; whereas other portion claim that treated surface is formed by a single-phase γ_N layer¹⁶,¹⁷,²⁰,²³,²⁴,²⁶. The two-phase structure formation would be due to the nitrogen enriching both the austenite and ferrite phases, but with no change on its respective stabilities, whereas the single-phase γ_N layer formation would be due to the nitrogen content enhancement leading the austenite phase stability to be attained also in ferrite grains. Regarding the treatment temperature limit values to prevent chromium nitride precipitation, some authors⁸ claim that the Cr-nitride phases precipitation occurs for nitriding temperatures from 400°C, others¹⁰ show that this only happens for nitriding temperatures above to 450 °C. One can note that the steel type, processing conditions, and the applied technique can lead to significantly different results regarding these aspects.

Table 1
Summary of the main contributions found in the open literature on low-temperature nitriding of duplex stainless steels.

References	Nitriding technique	Duplex stainless steel	Nitriding condition		Main results
			Time (h)	Temperature (°C)
Pinedo et al.7	d.c.	UNS S31803	20	400	Formation of a two-phase γN + αN layer with 3.8 and 4.4 wt.% N, respectively (αN thicker than γN, both with microhardness > 1350 HV).
Tschiptschin et al.8	d.c.	UNS S31803	20	400	Formation of a two-phase γN + αN layer, with 3.0 and 2.5 μm thickness, and 1360 HV and 1510 HV microhardness, respectively (intense coherent ε-Fe3N nitride precipitation inside αN).
Larisch et al.16	Gas nitriding	UNS S31803	3 to 12	400 to 550	Formation of a two-phase γN + αN layer formed in austenite and ferrite grains, respectively.
Calabokis et al.17	d.c. (LTPN)	UNS S32750	4	350, 400	Formation of a precipitation-free single-phase γN layer with significant corrosion resistance performance improvements for all treatment conditions.
Bielawski and Baranowska18	d.c.	UNS S31803	8 to 60	250 to 500	Formation of precipitation-free single-phase γN layer, with thickness of 15 μm and 8.9 wt.% N produced at temperatures up to 400 °C.
Kliauga and Pohl19	high-frequency pulsed plasma	UNS S31803	20 and 40	350 and 400	At 350 °C, the layer homogeneously covers the ferrite and the austenite phase regions, the latter phase showing γN formation, and needle-like α′′-Fe16N2 precipitation occurring in the layer-ferrite phase interface. At 400°C, Cr2N + Fe2-3N + γ'-Fe4N nitrides were present in the layer showing cracks.
Blawert et al.20	PI3	UNS S31803	3	400	The pre-existing ferrite phase transforms into γN phase, which in the extreme surface decomposes into ferrite + martensite + CrN.
Li et al.21	d.c.	UNS S82441	10	390 to 480	The original austenite phase become γN phase, and the ferrite phase is supersaturated with N forming ε-Fe3N nitride precipitates.
Chiu et al.22	ASPN	UNS S31803	10 to 25	400 to 450	At 420 °C, a two-phase γN and αN layer is formed, and at 435 and 450 °C, Fe and Cr nitrides precipitate in the layer.
Nagatsuka et al.23	ASPN and d.c.	UNS S32550	5	400 and 450	Single-phase γN layer is obtained (in austenite as well as in ferrite phases) on both d.c. and ASPN samples treated at 400 and 450 °C.
Alphonsa et al.24	d.c.	UNS S32205	4	350 to 500	GIXRD results show that at 350 and 400°C only γN phase is present in the layer and that at 450 and 500 °C Fe3N and CrN phases form.
Assmann et al.25	GDN and CIBNI	UNS S32101	3	300, 350 and 380	Α γN layer with nitride precipitation arising from 300 °C, with 1220 HV microhardness is obtained for the poor nitrogen atmosphere GD process. The CIBNI process even at high temperature produced only nitrogen-expanded austenite.
Blawert et al.26	PI3	UNS S31803	1, 3 and 7	200,400 and 500	Single-phase γN layer is formed on both austenite and ferrite grains.
Oliveira et al.27	PI3	UNS S32750	3	292 to 401	ε-Fe2-3N and γ-Fe4N iron nitrides are formed in the modified ferrite grains, whereas γN was produced mostly in austenite grains. Overall near surface microhardness is increased from 610 to 1530 HV.
Bobadilla and Tschiptschin28	d.c.	UNS S31803	4	350 to 500	The non-unidirectional N diffusion tends to result in a N diffusion flow from ferrite to austenite, leading to the formation of a duplex γN + αN layer with small thickness difference in both phases.

• d.c.: direct current; PI³: plasma immersion ion implantation; ASPN: active screen plasma nitriding; GDN: glow discharge nitriding; and CIBNI: conventional ion beam nitrogen implantation.

It is also to be noted that the UNS S32750 SDSS, studied here, presents a high 'Pitting Resistance Equivalent Number' (PREN = wt.% Cr + 3.3 × wt.% Mo + 20 × wt.% N), in the case higher than 40²⁹, which makes it a material indicated for extremely corrosive environments. As seen from the above-indicated PREN equation, by alloying the UNS S32750 SDSS surface with nitrogen better corrosion resistance should be expected, according to ref³.

Using a strong metallurgical background and the knowledge produced up to now on low-temperature nitriding of stainless steels, the aim of the present work was to conduct experiments to study the nitriding process kinetics of UNS S32750 SDSS substrates, bringing a deeper discussion on determining factors that influence the treatment on both steel phases. Considering that the chemical composition of the substrate plays important role on the layer formation in thermochemical treatments of different stainless steels³⁰,³¹, this same premise should be also valid for each phase of DSSs and SDSSs, which present different compositions and structures. So, this work reports an experimental investigation on the thermodynamic and kinetic view-point of the nitrided layer growth behavior, carried out separately at the ferrite and austenite phase grains for the UNS S32750 super duplex stainless steel (SDSS), as a function of the nitriding temperature and time.

2. Experimental Procedure

Samples of the UNS S32750 SDSS, with 20 × 20 × 3 mm³, were cut by wire electrical discharge machining from a 2.5-inch rolled steel rod. Semi-quantitative X-ray fluorescence analysis was used to determine the samples chemical composition, which showed 25.6% Cr, 6.8% Ni and 4.1% Mo (in wt.%). Samples were grinded and polished with 1 μm Al₂O₃ abrasive suspension and 1 μm diamond paste to obtain a mirror polished finishing. Before nitriding, samples were cleaned in ultrasonic bath using isopropyl alcohol, dried in a heated airflow and introduced into the discharge chamber. The treated surfaces were always those perpendicular to the rod rolling direction. The average microhardness values obtained in the ferrite and austenite grains for untreated samples were 325 and 270 HV0.01, respectively.

Plasma nitriding was carried out in the same plasma system previously illustrated in Anjos et al.³². The plasma apparatus consisted basically of a square-wave pulsed DC power supply operating at 4.16 kHz, and a stainless steel discharge chamber of 350 × 380 mm (diameter × height), attached to steel plates and sealed with O-rings at both ends. In such system, sample heating was reached by plasma species (ions and fast neutrals) bombardment, being the samples temperature measured by means of a chromel–alumel thermocouple (K–type of 1.5 mm diameter) inserted 8 mm depth into the sample holder (see Anjos et al.³²).

After evacuating the system to a residual pressure of 3 Pa (22 mTorr) using a double stage mechanical vacuum pump, samples were negatively biased at the voltage of 700 V in order to perform its (plasma) sputter-cleaning. This nitriding steps was performed at 300 °C, for 0.5 h, using a gas mixture constituted of 80% H₂+ 20% Ar, at a flow rate of 3.34 × 10^-6N.m³.s^-1 (200 sccm), under 400 Pa (3 Torr) pressure. Two treatment series were studied aiming to determine the influence of the temperature and time parameters on the nitriding kinetics:

1. i

   one varying the nitriding temperature at 300, 350, 400 and 450 °C for 4 h; and,

2. ii

   the other one varying the nitriding time for 2, 4 and 8 h, at a fixed temperature of 350 °C.

In both series the treatments were performed employing a gas mixture constituted by 70% N₂+ 20% H₂+ 10% Ar, and the other parameters being kept at the same values as those used in the cleaning step. It is important to clarify that the higher nitriding temperature (450 °C) and longer time (8 h) parameters were chosen taking care to avoid the “475°C tempering embrittlement” in the studied steel.

For the microstructural characterization of the treated surfaces, polished samples cross-sections were etched using Behara reagent (20 ml hydrochloric acid + 80 ml distilled water + 1 g potassium metabisulfite + 2 g ammonium bifluoride). Microstructures were observed using a scanning electron microscope (SEM), in the case a TSCAN VEGA 3 apparatus equipped with energy dispersive X-ray microprobe (EDS) to carry out elemental analyses. The phases of the obtained layers were identified by means of X-ray diffractometry (XRD), using a Shimadzu XRD 7000 diffractometer with CuKα radiation (Kα₁ → λ = 1.54056 Å and Kα₂ → λ = 1.54439 Å), in the Bragg-Brentano (for 2θ angle varying on the 30–90° range) XRD configuration, using a scanning speed of 2°/min. Hardness measurements were performed on the treated (top of the sample) and untreated (bottom of the sample – not exposed to plasma) surfaces using a Shimadzu microdurometer type HMV-2 T, applying a load of 10 gf (9.8 × 10− 2 N), for a peak-load contact of 15 s. To be representative, due to the relatively large distribution of the measurement values in different phases, the layer thickness and microhardness values were obtained from the mean of fifteen measurements taken for grains of each phase.

Finally, the activation energy for the nitrogen diffusion into the treated layer was determined by the Arrhenius plot for both austenite and ferrite phases, following the procedure described in ref²⁸,³³,³⁴. From the Arrhenius-type behavior for the nitriding thickness data, considering that the process is diffusion controlled (being the layer thickness proportional to (Dt)^1/2, where D is the diffusion coefficient and t is the diffusion time), the activation energy (Q_d) for nitrogen diffusion can be calculated by the slope of the Arrhenius plot according to Equation 1:

$$\ln \xi =A-\left(\frac{Q_d}{2R}\right)•\frac{1}{T}$$ (1)

where ξ is layer thickness (μm), R is the universal gas constant (8.31 J mol⁻¹K⁻¹), and T corresponds to absolute temperature (K).

3. Results and Discussion

3.1. Effect of the variation of the nitriding temperature

Cross-section micrographs of samples treated at 300, 350, 400 and 450 °C, for 4 h are presented in Figure 1ad, respectively (As an additional support optical micrographs are also presented in Figure A1a-d of Appendix A). A continuous layer formed from the nitrogen diffusion in both steel phases can be observed for all studied conditions. As expected, thicker layers are obtained for higher treatment temperatures (also see Figure 2). Note that the original austenite and ferrite grains as well as the obtained layers have different reactivity to the Behara reagent chemical etching. Two distinct region kinds can be easily observed, as follows (see Figure 1ad):

Figure 1
Cross-section SEM micrographs of samples treated at: (a) 300, (b) 350, (c) 400, and (d) 450 °C. Treatments carried out for 4 h, using a gas mixture composed of 70% N₂+ 20% H₂+ 10% Ar at a flow rate of 3.34×10⁻⁶Nm³s⁻¹, and pressure of 400 Pa.

Figure 2
Arrhenius plot for the plasma nitrided layer thickness formed on austenite and ferrite grains of the UNS S32750 SDSS. Treatments carried out for 4 h, using a gas mixture composed of 70% N₂+ 20% H₂+ 10% Ar at a flow rate of 3.34×10⁻⁶Nm³s⁻¹, and pressure of 400 Pa.

1. i

   a nitrogen-chemically-altered region at the effectively treated surface, where the obtained layer (of light gray-aspect) is more resistant to etching than the other regions; and,

2. ii

   a non-chemically-altered region, at the substrate bulk, which can be subdivided in two other regions clearly showing the original steel phases, namely the ferrite grains (of dark-aspect) much more susceptible to etching, and the austenite grains (of gray-aspect) slightly less susceptible to etching.

This result is especially important in the present work, since the revealed microstructure is strong enough to identify and distinguish at least three different phases in the studied samples. In this case, the continuous layer in the nitrogen-chemically-altered region seems to be very homogeneous when formed on both the ferrite grains and the austenite grains, presenting totally different aspect regarding the original steel phases (valid for samples treated at 300, 350 and 400 °C in Figure 1ac, respectively). This result suggests that the continuous layer obtained in the nitrogen-chemically-altered region is constituted of a unique phase. Differently, the aspect of the 450 °C treated layer in Figure 1d additionally suggests precipitation of Cr-nitride phase along the nitrided layer formed over the ferritic and austenitic grains, agreeing with the indicated in ref²⁰,³⁵. It is important to highlight that, unlike what was shown here, other authors¹⁸,²²,²⁴ reported the occurrence of Cr-nitride precipitation only in the ferritic grains, for similar treatment temperatures. Since CrN precipitation is incipient, CrN is not directly observed. The observed regions are sites depleted in chromium, due to CrN precipitation, which are less chemically resistant to the Behara etchant.

The Arrhenius plot for the thickness of the layer formed on each phase is depicted in Figure 2. These data were elaborated considering that nitrogen diffusion is the limiting factor for the layer growth, and so, the obtained activation energy is considered representative of nitrogen diffusion in the treated layer. Even if it is good approximation, it is to be noted that, strictly speaking, the obtained values are in fact and effective activation energies since, for example, the N surface concentration is not the same for the different temperatures. From Figure 2 results, it can be observed that the obtained layer on the ferrite phase is slightly thicker than that on the austenite phase. Despite the well-known relationship between diffusion and crystalline structure and orientation³⁶, this argument alone is not sufficient to explain this result, regarding the nitrogen diffusion in DSSs, as it will be presented ahead considering at least three other different factors. The observed data linearity in the plot suggests that the LTPN of the UNS S32750 SDSS is a diffusion-controlled process and that atomic diffusion is the limiting mass transfer mechanism for the treated layer growth. It also suggests that the secondary phase precipitation did not significantly influence the process activation energy. Some authors²⁸-³⁴ show a deviation from the linearity for high temperature treatment conditions, indicating that the layer growth is affected by the secondary phase precipitation, leading to the increase of the activation energy needed for the layer growth. Although Figure 1d indicates the treated layer sensitization occurrence, thus Cr-nitride precipitation, it is believed that it was only in its initial stage, since it did not significantly interfere on the nitriding process kinetics.

The activation energy for the nitrogen diffusion into the treated layer formed on the austenite and ferrite grains, estimated by the Arrhenius plot slope (Figure 2), was 115±3.2 and 120±5.4 kJmol^-1 (equivalent to 1.19±0.03 and 1.24±0.05 eV, and to 27.5±0.7 and 28.7±1.3 kcal/mol), respectively. This slight difference in the obtained activation energies gives additional support to the assumption that a single-phase layer has been formed on both the phases, being this slight difference probably related to specific characteristics of the treated layer on each substrate phase kind, as discussed ahead. These activation energies agree very well with those of Bobadilla and Tschiptschin²⁸, which obtained values of 117 and 124 kJmol^-1 for the nitrogen diffusion activation energies in UNS S31803 DSS austenite and ferrite grains, respectively, but, differently from here, they suggest the obtainment of a two-phase nitrided layer. Results published in the literature indicate nitrogen diffusion activation energy values of 111.4 kJmol^–1 for PIM 316L ASS³⁷ and 107 kJmol^–1 for AISI 304L ASS³⁸. For ferritic stainless steels, no activation energy data were found in the literature, but much earlier, a study presented an activation energy of 76 kJmol^–1 for a low-alloyed ferrite (α-Fe) steel³⁹.

Starting from the data determined by the Arrhenius plot extrapolation and considering Equation 1, the evolution of the nitrided layer as a function of the treatment temperature in the austenitic grains of the UNS S32750 SDSS can be described by Equation 2, whereas that for ferritic grains, through Equation 3:

$\ln \xi \left(K/\mu m\right)=2.49-\frac{115kJmo{l}^{-1}}{2R}•\frac{1}{T}$ (2)

$$\ln \xi \left(K/\mu m\right)=2.46-\frac{120kJmo{l}^{-1}}{2R}•\frac{1}{T}$$ (3)

The higher thickness of the nitrided layer on the ferrite phase, according to ref⁸,¹⁶,¹⁷,²⁷,³⁶,³⁷, is justified by the higher nitrogen diffusion coefficient in the ferrite than in the austenite structure. In the present work, this justification makes sense only for the initial nitriding moments, when the nitrogen enrichment was not sufficient to promote the ferrite-to-austenite transformation. This seems to be true up to the moment the nitrogen-supersaturated ferrite would transform to nitrogen-alloyed austenite, which is believed to occur in accordance with XRD results presented ahead (see Figure 3). Finally, from Figure 1ad it is also possible to verify a slight tendency to a more prominent layer growth at the original grain boundaries between ferrite and austenite phases. Since these boundaries are high diffusivity paths that presents lower diffusion activation energy⁴⁰, such sites would facilitate the nitrogen diffusion and can explain the observed tendency.

Figure 3
XRD patterns for untreated sample and samples treated at 300, 350, 400 and 450 °C. Treatments carried out for 4 h, using a gas mixture composed of 70% N₂+ 20% H₂+ 10% Ar at a flow rate of 3.34×10⁻⁶Nm³s⁻¹, and pressure of 400 Pa. (where: γ – austenite phase; α – ferrite phase; γ_N – nitrogen-expanded austenite phase; and γ'_N – secondary nitrogen-expanded austenite phase.

Going further in the discussion, it is worth explaining the meaning of the activation energy regarding the specific atomic migration event that would be operating in the present treatment. As known, the nitrogen atoms migration is based on the interstitial diffusion mechanism, but the involved activation energy can be dependent on the environment where such atoms are diffusing, thus on the layer composition. This point is especially important here, since the studied stainless steel substrate presents two phases, with different compositions depending on the grain (phase) kinds. As the thermochemical treatment is performed in paraequilibrium the substrate substitutional alloying elements will not significantly diffuse during treatment. So, even if an expanded-austenite treated layer is formed over both grain kinds, the composition of the treated layer built on the different phases will be different, as a consequence of the substitutional alloying elements partition occurred during the steel production (see factor i, below). As soon, although the activation energy values calculated here for each phase are similar, the hypothesis is supported that different phenomena are involved and promote the formation of the nitrided layer in each of these phases, as discussed below.

On the other hand, regarding the point lattice defects produced from the plasma-surface interaction, besides the gradient of the nitrogen interstitial, new supposedly produced vacancies would mainly work on the first atomic layers of the treated material. Since the vacancy density gradient tends to rapidly decrease into the substrate bulk, it would be important for the increment of the nitrogen solubility at the outermost surface of the substrate. It is to be remembered that in low-temperature (paraequilibrium) nitriding of stainless steels, the substitutional diffusion mechanism that is strongly dependent on the equilibrium vacancy density, would be negligible.

Therefore, the nitrogen diffusion in treated surfaces of DSSs, besides depending on each phase crystalline lattice structure, would also depend on at least four more factors:

1. i

   The partitioned alloying elements effect on the nitrogen diffusivity: Ni atoms partitioned in austenite grains, and Cr and Mo^a atoms preferentially partitioned in ferrite grains², influence the nitrogen interstitial atom diffusion kinetics, as indicated is ref³⁰,³¹. It has been shown that Ni in MSSs promotes the layer thickness reduction in paraequilibrium nitriding³⁰. More, the nitrogen solubility decreases as the Ni content on Fe alloys is increased⁴¹. So, one can assume that the nitrogen diffusion in the treated layer formed on the ferrite is faster than in austenite;

2. ii

   The alloying elements trapping effect: diffusion or trap sites are usual positions occupied by interstitial atoms diffusing into SSs at low temperature⁴². Nitrogen easily diffuses from diffusion sites due to its lower activation barrier (1.45 eV⁴³), but it has difficulty to diffuse from trap sites, showing greater activation barrier (2.5 eV⁴³), due to Cr atoms. Thus, as Cr as well as Mo atoms are preferentially in ferrite grains², it is assumed that a higher nitrogen diffusion activation energy in the treated layer on ferrite grains is due to the trapping effect;

3. iii

   The different stress generated in both phases of DSSs: during heating, in the initial stage of the samples treatment, different expansion coefficients of ferrite and austenite phases tend to result in compressive residual stresses in ferrite, while equilibrium tensile stresses are generated in the austenite⁴⁴. Tensile stresses promote a decrease on the activation energy for interstitial atoms diffusion⁴⁵, whereas compressive stress can retard its fast diffusion⁴⁶, opposing, at least in part, the factor i. So, it is assumed that at the beginning of the treatment nitrogen diffusion in the austenite phase, subjected to tensile stress, is faster than in ferrite grains that are subjected to compressive stress. Considering that the nitrogen-alloyed layer is constituted of expanded-austenite, the phase transformation energetic barrier in original ferrite grains must be overcome to initiate the layer growth; and

4. iv

   The stress resulting from ferrite-to-austenite transformation due to nitrogen supersaturation: After being supersaturated by nitrogen atoms, the ferrite structure in the treated layer is transformed in expanded-austenite (γ_N). The present factor is corroborated by results published in ref⁴⁷,⁴⁸. In Allenstein et al.⁴⁷, strong evidences of tempered martensite (or bcc structure ferrite)-to-nitrogen-expanded austenite transformation was observed in CA6NM MSS steel after LTPN. In Zangiabadi et al.⁴⁸, conclusive evidence for martensitic nature phase isothermal transformation from martensite/ferrite-to-austenite was observed at paraequilibrium temperature nitriding (LTN) of 15-5 PH (precipitation hardening), 17-7 PH, and 2205 (duplex) stainless steel surfaces with nitrogen concentrations from 8 to 16 at.%⁴⁸. Since the ferrite-to-austenite transformation occurs with volumetric contraction⁴⁹, an eventual tensile residual stress state at the transformed surface would contribute to a greater nitrogen diffusivity. It is to be remembered that the cause of the volume and packing density differences for both the considered iron (phases) crystalline structures is also related to other factors such as the coordination number, this one affecting the effective atomic (and of radii) ratios, besides the possible different spin electronic states, much more than simply on those results obtained from the hard sphere model.

From our results, it seems that factors iii and iv influence the nitrogen diffusion and alloying through the layer formed on the ferrite phase, leading its structure and composition to be changed along the treatment time and surely all over the treated layer depth, remembering that factors i and ii remain unchanged in this case, since the substitutional atoms diffusion is negligible in paraequilibrium¹⁹. Based on all above-presented aspects, even that factors ii and iii tend to promote an activation energy increase for the layer built in the ferrite phase, allied to the similar role promoted by the higher molybdenum atoms in this phase³¹,⁴⁵, as confirmed ahead (see Figure 4 and Table 2 results), the nitrogen-supersaturated ferrite-to-austenite transformation would explain the similar layer growth activation energy values obtained in our work on both phases as well as the similar treated layer growth rates observed in both phases after the referred transformation. This also should hold for all nitriding process of DSS where this transformation occurs and that is limited by the N atoms diffusion kinetics. This assumption is supported by observing treated layer thickness values indicated in Figure 2. Regarding the diffusion in austenite grains, its higher chromium and molybdenum contents in the UNS S32750 SDSS when compared to that found in the AISI 316³⁷ and AISI 304³⁸ ASSs would also lead the nitrogen diffusion activation energy in the SDSS austenite phase to be slightly higher than that in ASSs.

Figure 4
EDS nitrogen and nickel concentration line profile along the layer obtained for samples treated at: (a) 300, (b) 350, (c) 400, and (d) 450 °C. Treatments carried out for 4 h, using a gas mixture composed of 70% N₂+ 20% H₂+ 10% Ar at a flow rate of 3.34×10⁻⁶Nm³s⁻¹, and pressure of 400 Pa.

Table 2
EDS elemental analysis characterization of the nickel (austenite-stabilizer), and chromium and molybdenum (ferrite-stabilizer and stable nitride formers alloying elements) partition on non-treated UNS S32750 SDSS samples.

Phase	Ni / Cr / Mo (wt.%)	Ni / Cr / Mo (at.%)
Austenite	8.5 / 24 / 5.5	8.1 / 26 / 3.2
Ferrite	5.5 / 27 / 7.0	5.2 / 29 / 4.0

Figure 3 presents XRD patterns of untreated and treated samples. Patterns for untreated samples show only peaks referred to ferrite and austenite phases, as expected. XRD patterns of treated samples show the occurrence of nitrogen-expanded austenite (γ_N / γ'_N) phase peaks, thus strongly suggesting the single-phase expanded-austenite (γ_N) layer formation. Note that the γ_N phase peaks are asymmetric, probably due to the depth nitrogen compositional gradient, varying from surface into the substrate bulk. Broadened and reduced-intensity diffraction peaks due to crystalline lattice residual stresses, displaced to smaller 2θ angles, confirm the crystalline lattice expansion caused by the nitrogen atoms alloying, as indicated in ref¹⁷,⁵⁰. In addition to γ_N peaks, samples treated at 300 and 350 °C also show original substrate peaks for both austenite and ferrite phases, which are no longer observed at 400 and 450 °C. The presence of original substrate peaks would indicate that the XRD probed depth is greater than the obtained layer thickness⁵¹.

For all cases studied here, the obtainment of nitrogen-expanded ferrite cannot be assured from our XRD results (see Figure 3), giving additional support for the single-phase γ_N layer assumption, and agreeing with the notion that nitrogen is a strong austenite stabilizer and substitutional diffusion is effectively frozen at the paraequilibrium temperatures ⁴⁸. In addition, it is possible to verify the existence of less intense peaks near the original austenite (γ) phase peaks, which could be attributed to a secondary nitrogen-expanded austenite phase layer, which is indicated as γ'_N⁵². A double layer comprising a double-structure nitrogen-expanded austenite layer was also recently observed in Toscano et al.⁵³. According to Toscano et al.⁵³, the occurrence of this phase indicates the formation of nitrogen-expanded austenite layer with double structure, namely an outermost layer with a higher nitrogen content (varying on the 20-26 at.% N range) and an sublayer with lower nitrogen content (varying on the 4-10 at.% N). From Williamson et al.⁵⁴, it is stated that the γ'_N phase formation is caused by stress-assisted diffusion, but here it could be also related to the just transformed ferrite between the outermost γ_N phase layer and the non-chemically altered substrate bulk. The suggested ferrite-to-austenite transformation in the nitrogen-enriched layer in the present work could occur from the mechanism presented by Zangiabadi et al.⁴⁸ comprising the martensitic austenite formation, as well as by Forgas et al.⁵⁵, which proved this transformation to occur by reverse strain-induced martensitic transformation for a DSS. Being of interest of the reader, major details of these mechanisms can be found in ref⁴⁸,⁵²,⁵⁵-⁵⁸.

Finally, still from Figure 3 results, it is possible confirming that the higher nitriding temperature the greater is the γ_N peaks displacement to smaller angles, evidencing the higher nitrogen supersaturation of the obtained layer, which is in agreement with the EDS analysis presented in Figure 4ad. Considering the possible sensitization observed in Figure 1d (and also Figure A1d, suggesting Cr-nitride phase precipitation, the absence of nitride phase peaks in the XRD pattern of samples treated at 450 °C would be due to the supposed small volumetric fraction of the formed nitride phase, being not detected with the applied XRD parameters.

EDS nitrogen and nickel elemental analysis longitudinally performed in the mean depth of the layer formed at 300, 350, 400 and 450 °C on both steel phases, regarding the original ferrite and austenite grains, exhibited average 15, 16, 17 and 23 at.% N (4.3, 4.6, 4.9 and 7.0 wt.% N) contents, respectively (see Figure 4ad). Note that nickel content measurements agree well with the expected partition of this element in each phase, showing higher contents in austenite and smaller contents in the ferrite phase. In addition, results indicate the tendency of the nitrogen content in the layer to be higher in the layer formed over ferrite than in austenite phase, as mainly evidenced for samples treated at 400 and 450 C. This tendency is also related to the partition of the alloying elements in both steel phases. Higher amounts of chromium and molybdenum are evidenced in the ferrite grains, as presented in Table 2, which shows the nickel (austenite-stabilizer), and chromium and molybdenum (ferrite-stabilizer and stable nitride formers alloying elements) components partition on non-treated UNS S32750 SDSS samples. Using a variant of the Schaeffler diagram – the Kakhovskii constitutional diagram⁵⁹,⁶⁰ – from the Cr_eq (chromium-equivalent number) and Ni_eq (nickel-equivalent number) estimation for the studied steel, only 0.66 wt.% N in solid solution would be enough to ensure the austenite phase stability. This limit nitrogen content value seems to be surpassed by far to the ones observed here, which showed a 4.3 wt.% N minimum as above-indicated, also agreeing well with the suggested ferrite-to-austenite transformation.

Microhardness measurement results for samples treated at the different studied temperatures are shown in Figure 5. Measurements were carried out in the layer formed on both the austenite and ferrite grains at the sample top (surface exposed to plasma) and bottom (surface non-exposed to plasma, thus untreated). Results indicate microhardness increase from 406 to 1900 HV_0.01 on ferrite grains and from 385 to 1530 HV_0.01 on austenite grains. Such microhardness increment is due to the high nitrogen supersaturation and the consequent compressive residual stresses generated by the expansion of the crystalline lattice in both phases, promoting increasing superficial hardening effect for higher nitriding temperatures, as observed by Tschiptschin et al.⁸. This result corroborates those presented in Figure 3, in which it was evidenced an increase in the γ_N peaks displacement to smaller 2θ diffraction angles, indicating higher nitrogen supersaturation for higher treatment temperatures. In addition, the higher microhardness of the layer formed on ferrite phase is in agreement with its higher nitrogen concentration measured by EDS in this case, as shown in Figure 4ad. The greater microhardness observed in the γ_N layer formed on the ferrite grains is linked to its higher nitrogen content, since the higher partitioned chromium and molybdenum contents promote an intensified nitrogen attractiveness and trapping effect in the ferrite phase (see Table 2).

Figure 5
Surface microhardness for untreated sample and samples treated at 300, 350, 400 and 450 °C. Treatments carried out for 4 h, using a gas mixture composed of 70% N₂+ 20% H₂+ 10% Ar at a flow rate of 3.34×10⁻⁶Nm³s⁻¹, and pressure of 400 Pa.

Regarding the actual layer microhardness value, it is worth mentioning that the indentation depth must not exceed 20% of the treated layer thickness to ensure that the measured microhardness values are significant and minimally affected by the non-chemically altered substrate bulk properties, as indicated in Borcz et al.⁶¹. So, the estimated increase on the microhardness by comparing the top and bottom surface values, above-mentioned, is strongly dependent on the obtained layer thickness. Note that for samples treated at 300 and 350 °C, the microhardness test indentation depth varied on the 0.66 to 1.38 μm range, comparable to the obtained layer thickness, which was on the order of 0.65 and 1.85 μm for both austenite and ferrite phases (see Figure 2). Differently, at 400 and 450 °C, obtained layers were thick enough to avoid the influence of the softer substrate bulk on the measured layer microhardness, in the case 4.35 ± 0.1 μm and 9.15 ± 0.1 μm, respectively, on both phases (grains).

3.2. Effect of the variation of the nitriding time

Cross-section micrographs of samples treated at 350 °C for 2, 4 and 8 h are presented in Figure 6ac, respectively (As an additional support optical micrographs are also presented in Figure A2a-c of Appendix A). The observed microstructures are similar to those presented in Figure 1ac, with results also suggesting γ_N phase layer formation on both austenite and ferrite phases. EDS nitrogen and nickel elemental analysis performed along treated layers exhibited average 13, 16 and 18 at. % N (3.6, 4.6 and 5.3 wt.% N) contents for samples treated for 2, 4 and 8 h (see Figure 7ac, respectively). Taking into account the Ni_eq and Cr_eq values from the Kakhovskii constitutional diagram⁴⁷,⁵⁹, such nitrogen contents also result to an austenite structure. Regarding the variation of the nickel contents for different steel phases, results were similar to those obtained in Figure 4ad. For all studied treatment times, there is no evidence of sensitization in the obtained layers, indicating that the Cr nitride precipitation did not occur.

Figure 6
Cross-section SEM micrographs of samples treated at 350 °C for: (a) 2, (b) 4; and (c) 8 h. Treatments carried out using a gas mixture composed of 70% N₂+ 20% H₂+ 10% Ar at a flow rate of 3.34×10⁻⁶Nm³s⁻¹, and pressure of 400 Pa.

Figure 7
EDS nitrogen and nickel concentration line profile along the layer obtained for samples treated for: (a) 2, (b) 4, and (c) 8 h. Treatments carried out at 350 °C, using a gas mixture composed of 70% N₂+ 20% H₂+ 10% Ar at a flow rate of 3.34×10⁻⁶Nm³s⁻¹, and pressure of 400 Pa.

A plot of the layer thickness values as a function of the square root of the treatment time resulted in a linear relationship (Figure 8), confirming that the nitrogen transport into the treated steel surface is diffusion-controlled. Figure 8 analysis shows that, considering the standard deviation of the layer thickness values, lines can be fitted cross the origin point (abscissa x = 0, and ordinate y = 0)^b, suggesting that no incubation period is presented for the treated layer formation. From this result, it can be inferred that the adopted plasma sputter-cleaning stage was able to promote the complete native oxide layer removal, which acts as a barrier preventing the nitrogen diffusion into the steel surface. It also indicates that the treated layer formation/growth occurred in the initial treatment periods for both substrate phases. Still, from Figure 8, at the specific processing temperature of 350 °C, the nitrided layer growth on the austenitic and ferritic grains can be described, from the nitriding time – t, by the Equations 4 and 5, respectively:

Figure 8
Layer thickness as a function of the square root of the treatment time for layer formed on austenite and ferrite grains of the UNS 32750 SDSS. Treatments carried out at 350 °C, using a gas mixture composed of 70% N₂+ 20% H₂+ 10% Ar at a flow rate of 3.34×10⁻⁶Nm³s⁻¹, and pressure of 400 Pa.

$\ln \xi \left(h/\mu m\right)=0.995•t^{1/2}$ (4)

$\ln \xi \left(h/\mu m\right)=1.004•t^{1/2}$ (5)

Figure 9 shows XRD patterns obtained for samples treated at 350 °C for 2, 4 and 8 h. Results are similar to those obtained in the nitriding temperature series study, as previously shown in Figure 3. For all treatment times, original austenite phase peaks were also shifted to smaller 2θ angles, due to the high levels of nitrogen in solid solution, confirming the presence of nitrogen-expanded austenite (γ_N / γ'_N) phase layer, being that no evidence of nitrogen-expanded ferrite phase peaks was observed. Finally, in XRD patterns details, the tendency of the major γ_N phase peak to slightly shifting to greater angles with the increase of the treatment time was observed. This peak shift towards larger 2θ angles would suggest a possible relief of compressive residual stress in the treated material.

Figure 9
XRD patterns for untreated sample and samples treated at 350 °C for 2, 4 and 8 h. Treatments carried out using a gas mixture composed of 70% N₂+ 20% H₂+ 10% Ar at a flow rate of 3.34×10⁻⁶Nm³s⁻¹, and pressure of 400 Pa. (where: γ – austenite phase; α – ferrite phase; γ_N – nitrogen-expanded austenite phase; and γ'_N – secondary nitrogen-expanded austenite phase.

Figure 10 presents the microhardness measurements performed on treated and untreated surfaces. It is possible to note in both ferrite and austenite phases a slight layer microhardness increase as the nitriding time is increased. Layer microhardness values of 811, 893, and 968 HV_0.01 on ferrite grains, and 727, 775, and 833 HV_0.01 on austenite grains were obtained for 2, 4 and 8 h treated samples, respectively.

Figure 10
Surface microhardness for untreated sample and samples treated at 350 °C for 2, 4 and 8 h. Treatments carried out using a gas mixture composed of 70% N₂+ 20% H₂+ 10% Ar at a flow rate of 3.34×10⁻⁶Nm³s⁻¹, and pressure of 400 Pa.

4. Conclusions

The effect of the treatment temperature and time on the microstructure, microhardness, and growth kinetics of layers produced by plasma nitriding in paraequilibrium conditions on UNS 32750 super duplex stainless steel samples was investigated. A deepened discussion comparing literature works on the present subject was performed, aiming to bring lights on the layer growth mechanism on the steel original austenite and ferrite phases from our results. The main conclusions for the studied conditions can be listed as follows:

• Low-temperature plasma nitriding can be successfully applied to improve surface microhardness of UNS 32750 SDSS, which is due to the formation of a single-phase γ_N nitrided layer formed on both austenite and ferrite phase grains;

• For the treatment carried out at 450 °C microstructural results show significant evidences of incipient chromium nitride precipitation in the obtained layer and should be avoided since it will reduce the treated surface corrosion resistance, which is not acceptable for the case of stainless steels;

• The nitrogen diffusion in the studied SDSS surface, besides depending on each phase crystalline lattice structure, should also depend on at least four more factors: i) the partitioned alloying elements effect on each steel phase; ii) the distinct alloying elements trapping effect; iii) the generated (tensile in austenite, and compressive in ferrite) stress effect from thermal variations in the treatment beginning; and iv) the effect of the stresses arising from the transforming phases due to nitrogen supersaturation.

• Maximum microhardness of 1530 and 1900 HV_0.01 was obtained for the layer grown on austenite and ferrite phases, respectively.

• The nitrided layer growth kinetics is similar for both the austenite and ferrite steel phases, being compatible to that of a diffusion-controlled process. The calculated nitrogen diffusion activation energy was 115±3.2 kJmol^-1 and 120±5.4 kJmol^-1 for the treated layer formed on austenite and ferrite phase grains, respectively.

5. Acknowledgments

The authors wish to express thanks to the Laboratory of X-ray Optics and Instrumentation (LORXI-UFPR) and Electron Microscopy Center (CME-UFPR) for the use of the XRD and SEM equipment, respectively. This work was supported by CAPES - Coordenação de Aperfeiçoamento de Pessoal de Nível Superior - Brasil (Finance Code 001), Fundação Araucária do Estado do Paraná - Brasil (PRONEX - Grant N. 46744), CNPq - Conselho Nacional de Desenvolvimento Científico e Tecnológico – Brasil, CNPq-Universal Grant N. 482380/2012-8, and MCTI/CNPq/CT-Aquaviário Grant N. 456347/2013-5. Dr. C.J. Scheuer thanks to FAPERGS (Research Support Foundation of the Rio Grande do Sul State) by the financial support through Grant N. 19/2551-0001749-9.

• 1
  Molybdenum could also promote the trapping effect but, since chromium has superior affinity to nitrogen and is present in higher concentration in DSSs, authors suppose that its effect is mainly related to the lattice distortion.
• 2
  The straight line adjustment was performed by forcing its passage through the origin.

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Supplementary material

The following online material is available for this article:

Appendix A

Publication Dates

• Publication in this collection
  30 Mar 2022
• Date of issue
  2022

History

• Received
  14 Sept 2021
• Reviewed
  05 Mar 2022
• Accepted
  06 Mar 2022