**6. Diffusion bonding of FGSS**

In the solid-phase diffusion bonding process, the holding temperature (TH) plays a role to govern the microstructure. When TH is higher than the recrystallization temperature, the crystal grain easily grows to be coarse one within the bonding time and reduces the strength. Hence, TH must be lowered as possible to maintain high strength in the mechanical characteristics. In particular, this task in the diffusion bonding of AISI304 must also solve an issue to remove the passive film at low temperature. In the conventional process, TH becomes higher than 1100 K [28] to eliminate the oxide film on the stainless steel and to accelerate the bonding.

In recent years, new materials have been developed. **Figure 14** shows the newly developed full-martensitic stainless steel (hereafter called WC). This material can accelerate the diffusion bonding at low temperature by introducing a large amount of strain into austenitic stainless steel before bonding. The process of introducing distortion into the material was carried out as follows and as shown in **Figure 14** [29]. First, blocks of AISI304 were cut into 40 × 40 × 20 mm samples (**Figure 15(A)**). The samples were then compressed, cut, and rolled into a 10-mm-thick sheet, so that the equivalent strain in the compression direction was 90% at 573 K (less than

**Figure 14.** *Developed full-martensitic stainless steel AISI304 (WC).*

**Figure 15.** *Schematics of distortion-introduction process. Arrows show compression direction.*

the recrystallization temperature) (**Figure 15(B)** and **(C)**). These warm-rolled samples were then rolled to 1-mm thickness at room temperature to transform the microstructure into martensite from austenite (**Figure 15(D)**). To evaluate the effects of the presence or absence of martensite on diffusion bonding, a sample was warm-rolled at 573 K (thus containing deformed austenite) with a reduction of 99% (hereafter called sample W99), as depicted in **Figure 15(E)**. For the cross-tensile test, all the bonding specimens were cut from the 1-mm-thick sheet, and then their bonding areas were mirror polished. The bonding experiment was executed at punching diameter 5 mm and high vacuumed.

**Figure 16** shows the effects of strain-induced martensitic and strain for fracture load of bonded AISI304 seats on bonding temperature [30]. WC/WC bonded sample only at 973 K has a fracture load of 2 kN, while TH must be increased up to 1073 K to attain the same fracture load in case of SOL/SOL bonded sample. That is, WC/WC

**Figure 16.** *Effect of strain-induced martensite and strain for fracture load of bonded SUS304 seats on boning temperature.*

**15**

**Figure 17.**

*Grain growth in FGSS by heat treatment for 1.8 ks at 1073 K.*

*Integrated Manufacturing of Fine-Grained Stainless Steels for Industries and Medicals*

full-martensitic stainless steel bonding has a capacity to decrease TH by 100 K against conventional SOL/SOL sample. Furthermore, this TH for WC/WC bonded sample to attain 2 kN is reduced down to 20,123 K by using the fine-grained W99/W99 boding. This difference in TH by 50 K between WC/WC and W99/W99 suggests that the martensitic phase works effectively when diffusion bonding stainless steel.

In summary, the holding temperature in diffusion bonding is lowered by using FGSS sheets to successfully fabricate the bonded parts without grain size coarsening. This selection of FGSS sheets for joining is suitable to produce the precision parts with

In standard heat and surface treatments, the stainless steel parts and tools are subjected to high-holding temperature [31]. The first issue of engineering is a crystalline coarsening by the grain growth. As shown in **Figure 17**, FGSS is easy to be coarsened by heat treatment for 1.8 ks at 1073 K [32]. Even below 1000 K, the thermal distortion is issued to deteriorate the microstructure of FGSS; precipitation of carbides and nitrides is also worried to lower the fatigue strength and corrosion toughness. Various nitriding processes were compared to explain how to lower the holding temperature in [33]. The gas nitriding with use of ammonia required for higher temperature than 973 K for industrial surface treatment. The liquid nitriding with use of cyan solutions was applied to surface modification of automotive parts under the temperature of 700–900 K. Among those nitriding processes, the plasma nitriding has capacity to harden the surface layer in the thickness of 0.1–1 mm without significant loss of corrosion toughness of stainless steels under lower holding temperature. This processing is categorized into two regimes as shown in **Figure 18**; e.g., high-temperature nitriding above 673 K and low-temperature nitriding below 673 K as surveyed in [34] for

Above the master curves in each stainless steel, the nitrided surface layer is hardened by the fine CrN (chromium nitride) precipitation with large volume fraction [35]. After removal of fragile *γ*′-Fe4N precipitated layer (so-called white layer) and diffusion treatment, the nitrided layer with the nitrogen solubility of 0.1–0.2 mass% is utilized as a hardened protective layer of tool-steel and stainless steel parts and dies in commercial [36]. The plasma nitrided AISI316 at "A" just above its master curve in **Figure 18** is mainly hardened by the CrN precipitation. On the other hand, the inner nitriding process at "B" and "C" below the master curve in **Figure 18** is governed by the nitrogen solid solution or the nitrogen

*DOI: http://dx.doi.org/10.5772/intechopen.89754*

**7. Surface treatment of FGSS**

austenitic stainless steels.

high strength.

*Integrated Manufacturing of Fine-Grained Stainless Steels for Industries and Medicals DOI: http://dx.doi.org/10.5772/intechopen.89754*

full-martensitic stainless steel bonding has a capacity to decrease TH by 100 K against conventional SOL/SOL sample. Furthermore, this TH for WC/WC bonded sample to attain 2 kN is reduced down to 20,123 K by using the fine-grained W99/W99 boding. This difference in TH by 50 K between WC/WC and W99/W99 suggests that the martensitic phase works effectively when diffusion bonding stainless steel.

In summary, the holding temperature in diffusion bonding is lowered by using FGSS sheets to successfully fabricate the bonded parts without grain size coarsening. This selection of FGSS sheets for joining is suitable to produce the precision parts with high strength.

## **7. Surface treatment of FGSS**

*Engineering Steels and High Entropy-Alloys*

**14**

**Figure 16.**

**Figure 15.**

*Effect of strain-induced martensite and strain for fracture load of bonded SUS304 seats on boning temperature.*

the recrystallization temperature) (**Figure 15(B)** and **(C)**). These warm-rolled samples were then rolled to 1-mm thickness at room temperature to transform the microstructure into martensite from austenite (**Figure 15(D)**). To evaluate the effects of the presence or absence of martensite on diffusion bonding, a sample was warm-rolled at 573 K (thus containing deformed austenite) with a reduction of 99% (hereafter called sample W99), as depicted in **Figure 15(E)**. For the cross-tensile test, all the bonding specimens were cut from the 1-mm-thick sheet, and then their bonding areas were mirror polished. The bonding experiment was executed at

*Schematics of distortion-introduction process. Arrows show compression direction.*

**Figure 16** shows the effects of strain-induced martensitic and strain for fracture load of bonded AISI304 seats on bonding temperature [30]. WC/WC bonded sample only at 973 K has a fracture load of 2 kN, while TH must be increased up to 1073 K to attain the same fracture load in case of SOL/SOL bonded sample. That is, WC/WC

punching diameter 5 mm and high vacuumed.

In standard heat and surface treatments, the stainless steel parts and tools are subjected to high-holding temperature [31]. The first issue of engineering is a crystalline coarsening by the grain growth. As shown in **Figure 17**, FGSS is easy to be coarsened by heat treatment for 1.8 ks at 1073 K [32]. Even below 1000 K, the thermal distortion is issued to deteriorate the microstructure of FGSS; precipitation of carbides and nitrides is also worried to lower the fatigue strength and corrosion toughness. Various nitriding processes were compared to explain how to lower the holding temperature in [33]. The gas nitriding with use of ammonia required for higher temperature than 973 K for industrial surface treatment. The liquid nitriding with use of cyan solutions was applied to surface modification of automotive parts under the temperature of 700–900 K.

Among those nitriding processes, the plasma nitriding has capacity to harden the surface layer in the thickness of 0.1–1 mm without significant loss of corrosion toughness of stainless steels under lower holding temperature. This processing is categorized into two regimes as shown in **Figure 18**; e.g., high-temperature nitriding above 673 K and low-temperature nitriding below 673 K as surveyed in [34] for austenitic stainless steels.

Above the master curves in each stainless steel, the nitrided surface layer is hardened by the fine CrN (chromium nitride) precipitation with large volume fraction [35]. After removal of fragile *γ*′-Fe4N precipitated layer (so-called white layer) and diffusion treatment, the nitrided layer with the nitrogen solubility of 0.1–0.2 mass% is utilized as a hardened protective layer of tool-steel and stainless steel parts and dies in commercial [36]. The plasma nitrided AISI316 at "A" just above its master curve in **Figure 18** is mainly hardened by the CrN precipitation. On the other hand, the inner nitriding process at "B" and "C" below the master curve in **Figure 18** is governed by the nitrogen solid solution or the nitrogen

**Figure 17.** *Grain growth in FGSS by heat treatment for 1.8 ks at 1073 K.*

**Figure 18.** *Relationship of the holding temperature to the holding time in the plasma nitriding.*

alloying into AISI316 without formation of nitride precipitates [37]. This low temperature plasma nitriding is proved to be suitable to the surface treatment of FGSS with high hardness and strength and without thermal distortion.

#### **7.1 Plasma nitriding of fine-grained AISI316 at 623 K**

The fine-grained AISI316 (FGSS316) specimen was prepared for plasma nitriding at 623 K [38–41]. Its chemical compositions are: [C] = 0.08 mass%, [Si] = 1.00 mass%, [Mn] = 2.00 mass%, [P] < 0.045 mass%, [S] < 0.030 mass%, [Ni] = 12.0 mass%, [Cr] = 17.0 mass%, and [Mo] = 2.5 mass% for iron in balance. The sample surface was mirror polished and cleaned by the ultrasonic cleaner before plasma nitriding. **Figure 19(A)** depicts the high density plasma nitriding system with use of the hollow cathode device to intensify the density of nitrogen ions as well as the NH-radicals. The plasma-processing conditions are also summarized in **Figure 19(B)** in correspondence to "B" in **Figure 18**.

Microstructure and nitrogen mapping on the cross section of this nitrided FGSS316 specimen, describe the average nitrogen diffusion layer from the surface. As shown in **Figure 20(A)**, the nitriding front end locates at the depth of 40 μm from the surface. Although the crystal grain size below this nitriding front end remains the same as before nitriding, these grains are significantly refined in the nitrided layer. The nitrogen mapping in **Figure 20(B)** proves that high nitrogen content uniformly distributes in the nitrided layer from the surface to the depth


#### **Figure 19.**

*High density plasma nitriding system with use of the hollow cathode for low temperature plasma nitriding of FGSS316. (A) Schematic view of nitriding system, and (B) plasma nitriding conditions in the following experiments.*

**17**

**Figure 21.**

*Integrated Manufacturing of Fine-Grained Stainless Steels for Industries and Medicals*

of 40 μm. This implies that the thick nitrided layer with high nitrogen content and

*Cross-sectional image of the plasma nitrided FGSS316 for 14.4 ks at 623 K. (A) SEM image and (B) nitrogen* 

**Figure 21** compares the XRD diagrams of FGSS316 specimen surfaces before and after plasma nitriding at 623 K for 14.4 ks. XRD diagram of bare FGSS316 is charac-

Besides for these three peaks, no other peaks are detected even by the narrow scanned XRD in **Figure 21**; no chromium and iron nitrides are synthesized in this plasma nitriding. Nitrogen solute atoms do not react with the constituent atoms such as iron and chromium in FGSS316 but work as a constituent alloying element in the *γ*-lattice of FGSS316. The nitrogen solute atoms occupy the octahedral vacancy sites in the *γ*-lattices in FGSS316. This in situ nitrogen solute occupation with vacancy sites in the *γ*-lattices accompanies with the *γ*-lattice expansion in elasticity in **Figure 21**, and, characterizes the nitrogen supersaturation process in the low-temperature nitriding. Difference in the holding temperature between "A" and "B/C" in **Figure 18** reflects on the inner nitriding mechanism. High-temperature nitriding is driven by the nitrogen body-diffusion process with the precipitation reaction of diffusing nitrogen atom and constituent elements in FGSS316 to nitrides. Low-temperature nitriding is controlled by the nitrogen boundary-diffusion with the nitrogen

as well as *γ* (200)-peak at 2θ = 50.7o

to 40.5o

, respectively.

, and, *γ* (200)-peak, from

is identified as *α*′ (110).

refined grains is formed by the present surface treatment at 623 K.

, respectively. A new peak detected at 2θ = 43.65o

*XRD diagram of FGSS316 specimens before and after plasma nitriding at 623 K for 14.4 ks.*

*DOI: http://dx.doi.org/10.5772/intechopen.89754*

terized by *γ* (111)-peak at 2θ = 43.5o

50.7o

**Figure 20.**

to 46.6o

*mapping analyzed by EDX.*

After nitriding, this *γ* (111)-peak shifts from 43.5o

*Integrated Manufacturing of Fine-Grained Stainless Steels for Industries and Medicals DOI: http://dx.doi.org/10.5772/intechopen.89754*

**Figure 20.**

*Engineering Steels and High Entropy-Alloys*

alloying into AISI316 without formation of nitride precipitates [37]. This low temperature plasma nitriding is proved to be suitable to the surface treatment of

The fine-grained AISI316 (FGSS316) specimen was prepared for plasma nitriding at 623 K [38–41]. Its chemical compositions are: [C] = 0.08 mass%, [Si] = 1.00 mass%, [Mn] = 2.00 mass%, [P] < 0.045 mass%, [S] < 0.030 mass%, [Ni] = 12.0 mass%, [Cr] = 17.0 mass%, and [Mo] = 2.5 mass% for iron in balance. The sample surface was mirror polished and cleaned by the ultrasonic cleaner before plasma nitriding. **Figure 19(A)** depicts the high density plasma nitriding system with use of the hollow cathode device to intensify the density of nitrogen ions as well as the NH-radicals. The plasma-processing conditions are also summarized in

Microstructure and nitrogen mapping on the cross section of this nitrided FGSS316 specimen, describe the average nitrogen diffusion layer from the surface. As shown in **Figure 20(A)**, the nitriding front end locates at the depth of 40 μm from the surface. Although the crystal grain size below this nitriding front end remains the same as before nitriding, these grains are significantly refined in the nitrided layer. The nitrogen mapping in **Figure 20(B)** proves that high nitrogen content uniformly distributes in the nitrided layer from the surface to the depth

*High density plasma nitriding system with use of the hollow cathode for low temperature plasma nitriding of FGSS316. (A) Schematic view of nitriding system, and (B) plasma nitriding conditions in the following* 

FGSS with high hardness and strength and without thermal distortion.

*Relationship of the holding temperature to the holding time in the plasma nitriding.*

**7.1 Plasma nitriding of fine-grained AISI316 at 623 K**

**Figure 19(B)** in correspondence to "B" in **Figure 18**.

**16**

**Figure 19.**

**Figure 18.**

*experiments.*

*Cross-sectional image of the plasma nitrided FGSS316 for 14.4 ks at 623 K. (A) SEM image and (B) nitrogen mapping analyzed by EDX.*

of 40 μm. This implies that the thick nitrided layer with high nitrogen content and refined grains is formed by the present surface treatment at 623 K.

**Figure 21** compares the XRD diagrams of FGSS316 specimen surfaces before and after plasma nitriding at 623 K for 14.4 ks. XRD diagram of bare FGSS316 is characterized by *γ* (111)-peak at 2θ = 43.5o as well as *γ* (200)-peak at 2θ = 50.7o , respectively. After nitriding, this *γ* (111)-peak shifts from 43.5o to 40.5o , and, *γ* (200)-peak, from 50.7o to 46.6o , respectively. A new peak detected at 2θ = 43.65o is identified as *α*′ (110). Besides for these three peaks, no other peaks are detected even by the narrow scanned XRD in **Figure 21**; no chromium and iron nitrides are synthesized in this plasma nitriding. Nitrogen solute atoms do not react with the constituent atoms such as iron and chromium in FGSS316 but work as a constituent alloying element in the *γ*-lattice of FGSS316. The nitrogen solute atoms occupy the octahedral vacancy sites in the *γ*-lattices in FGSS316. This in situ nitrogen solute occupation with vacancy sites in the *γ*-lattices accompanies with the *γ*-lattice expansion in elasticity in **Figure 21**, and, characterizes the nitrogen supersaturation process in the low-temperature nitriding.

Difference in the holding temperature between "A" and "B/C" in **Figure 18** reflects on the inner nitriding mechanism. High-temperature nitriding is driven by the nitrogen body-diffusion process with the precipitation reaction of diffusing nitrogen atom and constituent elements in FGSS316 to nitrides. Low-temperature nitriding is controlled by the nitrogen boundary-diffusion with the nitrogen

**Figure 21.** *XRD diagram of FGSS316 specimens before and after plasma nitriding at 623 K for 14.4 ks.*

supersaturation in concurrent. This difference reflects on the microstructure evolution during the inner nitriding. Let us analyze this by using EBSD.

**Figure 22(A)** depicts the phase mapping on the cross section of nitrided FGSS316 at 623 K. The original microstructure below the nitriding front end has mainly austenitic phase but partially includes the martensitic phase. This phase transformation is strain-induced from the original *γ*-phase by intense rolling during fabrication of FGSS316 plates. Total volume fraction of *α*′-phase is only 6%. This implies that highly strained *γ*-phase grains must be massively transformed into *α*′ phase and that most of *γ*-phase remains as a matrix. On the other hand, this martensitic phase distributes finely together with *γ*-phase above the nitriding front end. The volume fraction of this *α*′-phase reaches to 70%. This implies that the nitrided layer has fine *γ*-*α*′ two-phase structure. The original matrix structure, seen below the nitriding front end in **Figure 22(A**), disappears and turns to be fine two-phase structure with nitrogen supersaturation.

As before mentioned, this nitrogen supersaturation accompanies with the plastic straining as well as the *γ* to *α*′ phase transformation by elastic straining in the above. **Figure 22(B)** shows the KAM distribution on the same cross section. Besides for the upper part with remaining relatively large *γ*-phase regions in **Figure 22(A)**, the two-phase structured grains in the nitrided layer have high-angled misorientation; almost every grain in the nitrided layer is plastically strained. As explained by **Figure 21**, the nitrogen supersaturated zones in each grain are forced to expand themselves in elasticity; while unsaturated zones have no elastic strains. Considering that strain incompatibility might be induced between these nitrogensaturated and unsaturated zones in every nitrided grain, the unsaturated zones are plastically strained. KAM distribution in **Figure 22(B)** represents the equivalent plastic strain distribution by this plastic straining.

**Figure 22(C)** shows the inverse pole figure in the nitrided layer. The original FGSS matrix, seen above the nitriding front end, turns to have much refined microstructure. The grain size is measured to be less than 100 nm or the spatial resolution limit of EBSD. This grain-size refinement is driven by intense plastic straining as seen in **Figure 22(B)**.

#### **Figure 22.**

*Microstructure on the cross section of nitrided FGSS316 at 623 K for 14.4 ks. (A) Phase mapping, (B) KAM distribution, and (C) inverse pole figures.*

**19**

**Figure 23.**

*(B) nitrogen solute content.*

*Integrated Manufacturing of Fine-Grained Stainless Steels for Industries and Medicals*

features by hardness testing and nitrogen depth profile measurement.

**Figure 23(A)** depicts the hardness depth profile of nitrided FGSS316 at 623 K. Hardness in the nitrided layer with the thickness of 40 μm is uniform by 1400 HV, and, gradually decreases down to the matrix hardness across the nitriding front end. The grain-size refinement in **Figure 22** contributes to this high hardness. As reported in [31, 33, 35], the maximum surface hardness by the high temperature plasma nitriding above 750 K was often limited by 1200–1300 HV. That hardness monotonously decreased with the depth except for high chromium alloys [31, 35]. Hence, the fact that nitrided layer is uniformly hardened layer by 1400 HV in average even without nitride precipitation, is a new knowledge to surface treatment of stainless steels and tool steels. As also stated in [31, 33, 35, 43], the maximum nitrogen solubility except for the bound nitrogen atoms in nitrides is limited by 0.1–0.2 mass%. Most of chromium solutes are reacted with diffusing nitrogen atoms and bound into CrN in the nitrided stainless steels over 750 K; those nitrided steels reduce their original corrosion toughness [44]. **Figure 23(B)** shows the nitrogen solute content depth profile. The surface nitrogen content becomes higher than 6 mass%, and, this [N] is nearly constant by 4–5 mass% down to the nitriding front end. As had been studied in [43, 45], this high nitrogen solute content significantly promotes the corrosion toughness even under the dipping condition into HCl solutions. This improvement of mechanical and electro-chemical properties by the

This low temperature plasma nitriding was first applied to fabricate the FGSS316 punch for microembossing the regular-square meshing pattern with the line width of 50 μm and pitch of 250 μm into aluminum plate. In the conventional die-fabrication, this textured multi-head punch is made by drilling each regularsquare microcavity with the edge of 200 μm. The number of microcavities reaches to 3200 on the punch with the size of 20 × 10 mm. When using the milling tool with the diameter of 10 μm and machining a single microcavity by the speed of 50 μm/s, including the cutting path changing time, and the threading depth of 10 μm, the

*Hardness and nitrogen solute content depth profiles of nitrided FGSS316 at 623 K. (A) Hardness and* 

This in situ formation of two-phase microstructure with refined grains in **Figure 22** is suitable to industrial applications. Different from the coarse mixture of *γ*-phase and *α*′-phase grains in normal two-phase stainless steels [42], this fine *γ*-*α*′ microstructure results in high hardness and toughness. In addition, high nitrogen solute concentration improves the corrosion toughness even in severe conditions. Let us evaluate on these

*DOI: http://dx.doi.org/10.5772/intechopen.89754*

**7.2 In situ two-phase structuring and refinement**

plasma nitriding is attractive to industrial applications.

**7.3 Plasma nitrided FGSS316 dies at 673 K**

*Integrated Manufacturing of Fine-Grained Stainless Steels for Industries and Medicals DOI: http://dx.doi.org/10.5772/intechopen.89754*

#### **7.2 In situ two-phase structuring and refinement**

*Engineering Steels and High Entropy-Alloys*

structure with nitrogen supersaturation.

plastic strain distribution by this plastic straining.

seen in **Figure 22(B)**.

supersaturation in concurrent. This difference reflects on the microstructure evolu-

As before mentioned, this nitrogen supersaturation accompanies with the plastic straining as well as the *γ* to *α*′ phase transformation by elastic straining in the above. **Figure 22(B)** shows the KAM distribution on the same cross section. Besides for the upper part with remaining relatively large *γ*-phase regions in **Figure 22(A)**, the two-phase structured grains in the nitrided layer have high-angled misorientation; almost every grain in the nitrided layer is plastically strained. As explained by **Figure 21**, the nitrogen supersaturated zones in each grain are forced to expand themselves in elasticity; while unsaturated zones have no elastic strains. Considering that strain incompatibility might be induced between these nitrogensaturated and unsaturated zones in every nitrided grain, the unsaturated zones are plastically strained. KAM distribution in **Figure 22(B)** represents the equivalent

**Figure 22(C)** shows the inverse pole figure in the nitrided layer. The original FGSS matrix, seen above the nitriding front end, turns to have much refined microstructure. The grain size is measured to be less than 100 nm or the spatial resolution limit of EBSD. This grain-size refinement is driven by intense plastic straining as

*Microstructure on the cross section of nitrided FGSS316 at 623 K for 14.4 ks. (A) Phase mapping, (B) KAM* 

**Figure 22(A)** depicts the phase mapping on the cross section of nitrided FGSS316 at 623 K. The original microstructure below the nitriding front end has mainly austenitic phase but partially includes the martensitic phase. This phase transformation is strain-induced from the original *γ*-phase by intense rolling during fabrication of FGSS316 plates. Total volume fraction of *α*′-phase is only 6%. This implies that highly strained *γ*-phase grains must be massively transformed into *α*′ phase and that most of *γ*-phase remains as a matrix. On the other hand, this martensitic phase distributes finely together with *γ*-phase above the nitriding front end. The volume fraction of this *α*′-phase reaches to 70%. This implies that the nitrided layer has fine *γ*-*α*′ two-phase structure. The original matrix structure, seen below the nitriding front end in **Figure 22(A**), disappears and turns to be fine two-phase

tion during the inner nitriding. Let us analyze this by using EBSD.

**18**

**Figure 22.**

*distribution, and (C) inverse pole figures.*

This in situ formation of two-phase microstructure with refined grains in **Figure 22** is suitable to industrial applications. Different from the coarse mixture of *γ*-phase and *α*′-phase grains in normal two-phase stainless steels [42], this fine *γ*-*α*′ microstructure results in high hardness and toughness. In addition, high nitrogen solute concentration improves the corrosion toughness even in severe conditions. Let us evaluate on these features by hardness testing and nitrogen depth profile measurement.

**Figure 23(A)** depicts the hardness depth profile of nitrided FGSS316 at 623 K. Hardness in the nitrided layer with the thickness of 40 μm is uniform by 1400 HV, and, gradually decreases down to the matrix hardness across the nitriding front end. The grain-size refinement in **Figure 22** contributes to this high hardness. As reported in [31, 33, 35], the maximum surface hardness by the high temperature plasma nitriding above 750 K was often limited by 1200–1300 HV. That hardness monotonously decreased with the depth except for high chromium alloys [31, 35]. Hence, the fact that nitrided layer is uniformly hardened layer by 1400 HV in average even without nitride precipitation, is a new knowledge to surface treatment of stainless steels and tool steels. As also stated in [31, 33, 35, 43], the maximum nitrogen solubility except for the bound nitrogen atoms in nitrides is limited by 0.1–0.2 mass%. Most of chromium solutes are reacted with diffusing nitrogen atoms and bound into CrN in the nitrided stainless steels over 750 K; those nitrided steels reduce their original corrosion toughness [44]. **Figure 23(B)** shows the nitrogen solute content depth profile. The surface nitrogen content becomes higher than 6 mass%, and, this [N] is nearly constant by 4–5 mass% down to the nitriding front end. As had been studied in [43, 45], this high nitrogen solute content significantly promotes the corrosion toughness even under the dipping condition into HCl solutions. This improvement of mechanical and electro-chemical properties by the plasma nitriding is attractive to industrial applications.

### **7.3 Plasma nitrided FGSS316 dies at 673 K**

This low temperature plasma nitriding was first applied to fabricate the FGSS316 punch for microembossing the regular-square meshing pattern with the line width of 50 μm and pitch of 250 μm into aluminum plate. In the conventional die-fabrication, this textured multi-head punch is made by drilling each regularsquare microcavity with the edge of 200 μm. The number of microcavities reaches to 3200 on the punch with the size of 20 × 10 mm. When using the milling tool with the diameter of 10 μm and machining a single microcavity by the speed of 50 μm/s, including the cutting path changing time, and the threading depth of 10 μm, the

#### **Figure 23.**

*Hardness and nitrogen solute content depth profiles of nitrided FGSS316 at 623 K. (A) Hardness and (B) nitrogen solute content.*

total machining time is estimated to be 1, 280 ks, or 356 h. This micromachining provides no solution to fabrication of the multi-arrayed punch with regular-square meshing heads. After [46, 47], the plasma printing method with use of the low temperature plasma nitriding provides a solution to fabricate this die as shown in **Figure 24**. A meshing pattern is first printed by using the screen with a unit pattern in **Figure 24(A)**. The unprinted cross-meshing surfaces are selectively nitrided at 673 K for 14.4 ks ("C" in **Figure 18**) in **Figure 24(B)**. After sand-blasting, this unit pattern transforms to **Figure 24(C)**. The microtexture arrayed punch is fabricated as shown in **Figure 24(D)** to have regular meshing texture in **Figure 24(E)**.

The takt time of this processing is only 18 ks, 70 times shorter than micromilling even excluding the cutting tool life as well as the preparation for CAM data before actual machining.
